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3D hierarchically porous Cu–BiOCl nanocomposite films: one-step electrochemical synthesis, structural characterization and nanomechanical and photoluminescent properties† ´,b Bradley J. Nelson,b Maria Dolors Baro ´, a Miguel Guerrero,*a Salvador Pane c d a ` nica Rolda ´n, Jordi Sort* and Eva Pellicer* Mo Three-dimensional (3D) hierarchically porous composite Cu–BiOCl films have been prepared by a facile onestep galvanostatic electrodeposition process from acidic electrolytic solutions containing Cu(II) and Bi(III) chloride salts and Triton X-100. The films show spherical, micron-sized pores that spread over the whole film thickness. In turn, the pore walls are made of randomly packed BiOCl nanoplates that are assembled leaving micro–nanopore voids beneath. It is believed that Cu grows within the interstitial spaces between the hydrogen bubbles produced from the reduction of H+ ions. Then, the BiOCl sheets accommodate in the porous network defined by the Cu building blocks. The presence of Cu tends to enhance the mechanical stability of the composite material. The resulting porous Cu–BiOCl films exhibit
Received 8th July 2013 Accepted 8th October 2013
homogeneous and stable-in-time photoluminescent response arising from the BiOCl component that spreads over the entire 3D porous structure, as demonstrated by confocal scanning laser microscopy. A broad-band emission covering the entire visible range, in the wavelength interval 450–750 nm, is
DOI: 10.1039/c3nr03491g
obtained. The present work paves the way for the facile and controlled preparation of a new
www.rsc.org/nanoscale
generation of photoluminescent membranes.
1.
Introduction
Three-dimensional (3D) porous lms have become increasingly popular in the last few years as they offer a number of applications in electrocatalysis,1 separation and concentration of gas molecules,2 batteries,3 sensors4 and electronic devices,5 among others. Several strategies are currently being pursued for the preparation of porous lms like de-alloying,6 galvanic replacement,7 or templating (e.g. using colloidal crystals as templates).8 Electrodeposition has recently emerged as a powerful and cheap technique to prepare 3D porous lms of metals and metallic alloys through different strategies: deposition in rigid porous templates (e.g. polystyrene opal templates),9 using hydrogen bubbles as a dynamic template10 or from electrolyte solutions containing suitable surfactants.11 The utilization of
Departament de F´ısica, Facultat de Ci`encies, Universitat Aut`onoma de Barcelona, E-08193 Bellaterra, Spain. E-mail: [email protected]; [email protected]
a
b
Institute of Robotics and Intelligent Systems (IRIS), Swiss Federal Institute of Technology (ETH) Zurich, CH-8092 Zurich, Switzerland
c
Servei de Microsc` opia, Universitat Aut`onoma de Barcelona, E-08193 Bellaterra, Spain
Instituci´ o Catalana de Recerca i Estudis Avançats (ICREA), Departament de F´ısica, Universitat Aut` onoma de Barcelona, E-08193 Bellaterra, Spain. E-mail: Jordi.Sort@ uab.cat
d
† Electronic supplementary 10.1039/c3nr03491g
information
12542 | Nanoscale, 2013, 5, 12542–12550
(ESI)
available.
See
DOI:
hydrogen co-evolution as a source of porosity during electrodeposition is a very convenient and fast way to prepare network porous structures of a number of metals (e.g. Cu,12 Ni,13 Pd,14 Pt,15 bimetallic Cu/Pd16) and alloys (e.g. Cu6Sn5).17 Among them, Cu is perhaps one of the most studied. Different electrolytes have been tested for the preparation of porous Cu lms with distinct pore morphologies, pore sizes and interconnectivities. Macroporous Cu lms constructed by nanodendritic walls have been obtained from acid sulphate solutions containing cetyltrimethylammonium bromide (CTAB).12 Cu foams with 3D interconnected spherical pore networks have been prepared in acid sulphate solutions containing HCl, polyethylene glycol (PEG) and 3-mercapto-1-propane sulfonic acid (MPSA) as additives.18 Remarkably, very few studies have reported the utilization of the hydrogen bubble template approach for the preparation of porous composite lms. A smart combination of metals with other compounds in a single 3D porous network structure could signicantly broaden the range of applications of porous lms. Bismuth oxychloride (BiOCl), a V–VI–VII group wide band gap semiconductor, is considered one of the most important bismuth oxyhalides. BiOCl has attracted an enormous amount of interest within the scientic community because of its excellent catalytic, optical, electric, magnetic and luminescent properties. As a result, BiOCl has found numerous uses as a
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Paper photocatalyst,19–21 or ferroelectric material,22,23 for many years. Moreover, it is widely used as a pigment in cosmetics, buttons and jewellery because it confers a semi-transparent appearance to the objects. Until now, several strategies have been reported for the preparation of BiOCl micro/nanostructures mostly in the form of powders: electrospinning,24 hydrothermal approach,25 reverse microemulsions26 or sonochemistry.27 However, for many industrial applications immobilized BiOCl lms are more adequate than powders. For example, problems encountered in nanopowder photocatalysts such as catalyst recovering (once the solution has been decontaminated) and particle aggregation can be circumvented if the active material can be properly immobilized without compromising its specic surface area. Nevertheless, studies on the preparation of nanostructured BiOCl lms are still scarce compared to the vast literature on BiOCl nanopowders. Cao and co-workers fabricated a BiOCl lm with ower-like hierarchical structure by reacting a Bi lm with H2O2/HCl solution.28 Very recently the preparation of BiOCl lms by electrodeposition of Bi and subsequent anodic oxidation has been reported.29 In all these studies the BiOCl lm is used for photocatalytic purposes. Other applications would also benet from having immobilized BiOCl lms like electrochemical sensing or optics. Herein, we present a facile, efficient and environmentally friendly approach for fabricating three-dimensional (3D) porous composite lms made of Cu and BiOCl components in one-step cathodic electrodeposition. An optimized bath formulation containing suitable concentrations of bismuth(III) and copper(II) chloride salts and Triton X-100 as an additive yields a crack-free porous composite structure. Also, the mechanical properties of the composite lms have been studied by nanoindentation in order to correlate the obtained data with both the porosity degree and the presence of Cu in the lms. Finally, the hierarchically porous Cu–BiOCl lms exhibit a strong and stable photoluminescence response in the entire visible range.
2.
Experimental section
2.1
Preparation of Cu-free and Cu–BiOCl composite lms
Cu-free and Cu–BiOCl composites were prepared by direct current electrodeposition (ED) in a thermostatized onecompartment three-electrode cell using a PGSTAT302N Autolab potentiostat/galvanostat (Ecochemie). The working electrode was positioned vertically within the electrolyte and consisted of 5 mm 6 mm Si chips (crystal orientation 100), on top of which a Ti adhesion layer of 50 nm and a Au seed-layer of 500 nm had been successively deposited through e-beam evaporation. The working area was 0.25 cm2. A double junction AgrAgCl (E ¼ +0.210 V/SHE) reference electrode (Metrohm AG) was used with 3 M potassium chloride (KCl) inner solution and an interchangeable outer solution. The outer solution was made of 1 M sodium chloride. A platinum sheet served as the counter electrode. For Cu–BiOCl composite electrodeposition, the electrolyte contained 1.2 M HCl, 8 103 M BiCl3, 8 103 M CuCl2, 2 102 M NaCl, and variable amounts of Triton X-100, saccharine and CTAB. The Cu-free BiOCl lms were prepared This journal is ª The Royal Society of Chemistry 2013
Nanoscale from an electrolyte free of CuCl2 salt. Analytical grade reagents and ultrapure water (18 MU cm) were used to prepare the electrolyte. The cell was lled with 50 mL electrolyte. All the samples were obtained galvanostatically by applying a constant current density of 1 A cm2. Deposition was conducted at room temperature under vigorous stirring (5000 rpm) using a magnetic stirrer bar. A similar electrical charge owed in all experiments, yielding lm thicknesses between 75 and 100 mm. Prior to deposition, the Au surface was degreased with acetone and dried with a so tissue. The Si backside was insulated by painting it with a nonconductive ink to ensure only the Au surface was conductive. Aer deposition, the coated substrates were rst dipped in 0.12 M HCl, aerwards dipped in Milli-Q water, and nally dried in air. 2.2
Morphological and structural characterization
The morphology and structure of the deposits was studied by scanning electron microscopy (SEM), high resolution transmission electron microscopy (HRTEM) and X-ray diffraction (XRD). SEM imaging was performed on a Zeiss Merlin microscope. TEM characterisation was carried out on a JEOL JEM2011 microscope operated at 200 kV. For TEM observations, a small area of sample surface was scratched-off and the powder released was dispersed in ethanol. Aerwards, one or two drops of the suspension were placed dropwise onto a holey carboncoated Cu TEM grid. The chemical composition of the lms was determined by energy dispersive X-ray spectroscopy (EDX). The pore size distributions were determined by manual analysis of enlarged micrographs by measuring ca. 100 pores on a given grid to obtain a statistical size distribution and a mean pore diameter. XRD patterns were recorded on a Philips X'Pert diffractometer using Cu Ka radiation, in the 20–90 2q range (0.03 step size, 8 s holding time). The global structural parameters, such as phase percentages, crystallite sizes, hDi (dened here as the average coherently diffracting domain sizes), and microstrains or atomic level deformations, h32i1/2, were evaluated by tting the full XRD patterns using the “Materials Analysis Using Diffraction” (MAUD) Rietveld renement program.30–32 Fourier transform infrared spectroscopy (FT-IR) was carried out on a Bruker TENSOR27 instrument. 2.3
Mechanical characterization
The mechanical properties (hardness, reduced elastic modulus and elastic recovery) were evaluated by means of nanoindentation, using a UMIS device from Fischer-Cripps Laboratories equipped with a Berkovich pyramidal-shaped diamond tip, operating in the load control mode. The maximum value of applied load was 0.6 mN to ensure that the lateral size of the indentation imprint was much smaller than the lm thickness. The thermal dri during nanoindentation was lower than 0.05 nm s1. Proper corrections for the contact area (calibrated with a fused quartz specimen), instrument compliance, and initial penetration depth (a pre-contact load of 0.2 mN was used) were applied. The hardness (H) and reduced elastic modulus (Er) values were derived from the load–displacement
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curves using the method of Oliver and Pharr.33 From the initial unloading slope, the contact stiffness, S, was determined as: S¼
dP dh
(1)
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where P and h denote the applied force and the penetration depth during nanoindentation, respectively. The reduced Young's modulus was evaluated based on its relationship with the contact area, A, and contact stiffness: pffiffiffiffi 2 S ¼ b pffiffiffiffi Er A p
(2)
where b is King's factor that depends on the geometry of the indenter (b ¼ 1.034 for a Berkovich indenter).34 The reduced modulus takes into account the elastic displacements that occur in both the specimen, with Young's modulus E and Poisson's ratio n, and the diamond indenter, with elastic constants Ei and ni (for diamond, Ei ¼ 1140 GPa and ni ¼ 0.07): 1 1 n2 1 ni 2 þ ¼ Er E Ei
(3)
The hardness was calculated from the following expression: PMax H¼ A
(4)
where PMax is the maximum load applied during nanoindentation. Finally, the elastic recovery was evaluated as the ratio between the elastic and total (plastic + elastic) energies during nanoindentation, Wel/Wtot. The values of Wel were calculated as the area between the unloading curve and the displacement axis. In turn, Wtot is the area between the loading curve and the displacement axis.34,35 To extract semi-quantitative information on the porosity level of the electrodeposited lms, the mechanical properties of non-porous bulk commercial BiOCl (Sigma-Aldrich, 98% purity) were also measured by nanoindentation. The results presented in this work are the statistical average of a set of 75 indentations for each sample. 2.4
Photoluminescence measurements
The photoluminescence (PL) properties of the lms were studied by Confocal Scanning Laser Microscopy (CSLM) analysis. Samples were mounted on Ibidi culture dishes (Ibidi GmbH, Martinsried, Germany) and imaged with a TCS-SP5 CSLM microscope (Leica Microsystems CMS GmbH Mannheim, Germany) using a Plan Apochromat 10/0.4 and 20/0.7 (dry) objectives. Both Cu-free and Cu–BiOCl composites were excited with an argon laser (488 nm) and the autouorescence was detected in the 500–785 nm range. The three-dimensional (3D) structure of lms was imaged using the same confocal microscope in reection mode with a 488 nm argon laser and detected in the 480–495 nm range. To determine the 3D microstructure, stacks of 50 to 80 sections were collected every 1 mm along the material's thickness. Three-dimensional models were generated from the xyz series using the Imaris X64 v. 6.2.0 soware (Bitplane; Z¨ urich, Switzerland). A series of images (xyl) were taken to determine the spectral emissions of the samples and to establish their maxima.
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Intrinsic uorescence was excited with a 405 nm line of a blue diode laser. Fluorescence emission was collected with a hybrid detector in 20 nm bandwidth increments (lambda step size ¼ 7 nm) in the range from 450 nm to 750 nm. A set of 30 regions of interest (ROIs) of 100 mm2 was used to analyze the mean uorescence intensity (MFI) of the samples in relation to the emission wavelength. Finally, photostability experiments were performed to monitor long time-lapse experiments. A 50 mW blue diode 405 nm laser was used at a fairly high output (Acousto-Optic Tunable Filter, AOTF ¼ 90%) in a eld of 1.5 mm 1.5 mm and the uorescence signal was detected in the range from 450 nm to 750 nm. Images were taken every 1 s for 10 min. A set of 20 ROIs of 100 mm2 were selected to show the MFI in the region in relation to time. Data from all studies were analyzed using the LAS AF soware ver. 2.4.1. (Leica Microsystems).
3.
Results and discussion
3.1 Synthesis, morphological and structural characterization of Cu-free and Cu–BiOCl composite lms Representative on-top SEM images of the Cu–BiOCl composite lms grown in the electrolyte containing 4 104 M of Triton X-100 are shown in Fig. 1. At low magnication, a uniform distribution of spherical pores of 5–10 mm in diameter (macroporosity) is observed over the entire surface of the lm (Fig. 1a and b). Moreover, the pore walls are not dense but inherently porous, with voids having dimensions from tenths of nanometers to a few microns (Fig. 1c). A close-up observation of the pore walls (Fig. 1d) reveals that they are in fact built from numerous randomly packed nanoplates, and that the nanoplates are stacked together either through plane-to-plane or edge-to-edge conjunctions, giving a ower-like appearance. Hence, the lms are hierarchically porous. The porosity spreads over the whole thickness as demonstrated by the cross-section image of the composite layer (Fig. 2). It is worth mentioning that the Triton X-100 additive suppresses the formation of cracks. With other additives like saccharine and CTAB cracked porous lms were obtained (ESI, Fig. S1†). As a result, the samples were so powdery and brittle that good contact between
Fig. 1 On-top SEM images of the Cu–BiOCl composite obtained in the electrolyte with 4 104 M Triton X-100 at different magnifications.
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Fig. 2 Cross-section SEM image of the Cu–BiOCl composite obtained in the electrolyte with 4 104 M Triton X-100.
the samples and the indenter tip could not be attained and the nanoindentation data became noisy and unreliable. For this reason, samples prepared using saccharine or CTAB as additives were not considered for in-depth mechanical and photoluminescence studies. Saccharine is typically used as a stressreliever and grain-rener in electrodeposited metals and alloys,35 but does not work for all systems and it has little effect on porosity. CTAB is a cationic surfactant with a positively charged head group that interacts electrostatically with the electrode surface. It gets adsorbed on the electroactive sites and, as a consequence, changes crystal growth mode and, in turn, grain morphology and size.36,37 Triton X-100 is recognized as an excellent wetting agent, which means that it favours H2 bubble detachment from the substrate,38 which likely promotes the evolution of porosity in the coating. At the same time, it would inhibit the formation of cracks in the lm as it builds. It is believed that such a synergetic effect is not achieved with saccharine and CTAB additives. Signicantly, Cu-free BiOCl lms exhibit different pore architecture (Fig. 3a). The lm surface is still very rough but the spherical macropores are no longer evident (Fig. 3b). A rose-like morphology is observed at high magnication (Fig. 3c). This suggests that the reduction of H+ (R1) that accompanies Cu discharge (R2) is responsible for the formation of the spherical macropores. The BiOCl nanoplates are then accommodated within the macropore walls dened by the Cu building blocks. The electrochemical reactions involved in the formation of the macropore skeleton are thus: H+ + 1e / ½H2
(R1)
Cu2+ + 2e / Cu
(R2)
On the other hand, the formation of BiOCl nanoplates would mainly proceed via the reaction of Bi(III) with an electrogenerated base (OH) in the presence of Cl (R3): Bi3+ + Cl + 2OH / BiOCl + H2O
(R3)
The formation of the BiOCl component does not take place by simply dipping the substrate in the electrolyte if no current is owing through the system (i.e., by hydrolysis). Hence, the formation of the BiOCl nanoplates necessarily involves a reaction with any species generated during the electrolysis (could be
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Fig. 3
(a–c) On-top SEM images of the Cu-free BiOCl film.
H2O2 that reacts with freshly reduced Bi, as well). Moreover, the electroreduction of Bi(III) toward metallic Bi would also take place: Bi3+ + 3e / Bi
(R4)
It has been reported that metallic bismuth electrodeposited at hydrogen evolution potentials exhibits a non-hierarchical porous dendritic structure.39 The potential–time (E–t) transients recorded during Cu-free and Cu–BiOCl lms are shown in the ESI, Fig. S2.† The E–t curves show an important decrease of the deposition potential toward more positive values during the rst few seconds (nucleation events) until stabilization around 3 V vs. Ag/AgCl is achieved. The stabilized potential is similar for both Cu-free and Cu–BiOCl lms and there is not any signicant decrease/increase of the average potential with time, which suggests that the surface area of the deposition front remains almost constant. The uctuations are due to the intense hydrogen evolution. Both the macropore size and size distribution of the Cu–BiOCl hierarchically layered composites can be greatly tuned by varying the Triton X-100 concentration in the electrolyte (ESI, Fig. S3†). Namely, both values decrease as the concentration of Triton X-100 is increased from 1 104 M (mean macropore size 24.7 mm) to 4 104 M (mean macropore size 9.1 mm). Energy-dispersive X-ray spectroscopy (EDX) analysis shows the presence of Bi, O and Cl (coming from the BiOCl layer) as well as Cu homogeneously distributed within the nanocomposite (Fig. 4).
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Fig. 4
Paper
EDX mapping distribution of Cu, Bi, O and Cl elements in the Cu–BiOCl composite film.
evidencing the nanocrystalline character of the composite electrodeposited lms. In turn, the values of microstrains (which are related to the density of dislocations and other structural defects) are relatively low, much lower than those typically obtained in heavily deformed metals,40 but still higher than the microstrain values in some electrodeposited nanocrystalline metallic alloys.35 To gain further information on Cu–BiOCl hierarchically porous lms, the product was further investigated by HRTEM. Fig. 6a shows a blend of nanosheets corresponding to BiOCl. The size of the sheets is well in the sub-100 nm range, in agreement with XRD results. Fig. 6b is a lattice-resolved HRTEM image taken on a single nanosheet. The corresponding lattice fringes have an interplanar distance of 0.27 nm, which matches the d-spacing of the (110) reection of tetragonal BiOCl (d110 ¼ 0.2753 nm). This is in accordance with the selected-area electron diffraction (SAED), indexed to a pure tetragonal phase (inset of Fig. 6a). This result indicates that the preponderant growth direction of the BiOCl layer is the [001] orientation, which is parallel to the (110) orientation. Metallic Bi on the BiOCl nanosheets is also detected, in agreement with XRD data (Fig. 6c). The inset shows the FFT of the area enclosed in red that belongs to the rhombohedral Bi phase (JCPDS no. 44-1246).39 The IR spectra of the sample show the typical bands of BiOCl in three different regions (2854 cm1, 1463 cm1 and 612 cm1) (ESI, Fig. S4†), again conrming the presence of the BiOCl component.
The crystallographic structure of the as-prepared Cu-free and Cu–BiOCl coatings was studied using XRD (Fig. 5). Apart from the reections coming out from the substrates (Si and Au), some of the diffraction peaks can be assigned to the tetragonal BiOCl phase (S.G., P4/nmm; JCPDS no. 06-0249, a ¼ b ¼ 0.3891 nm and c ¼ 0.7369 nm). In addition, a narrower reection at 2q ¼ 43.30 is identied in the XRD pattern of the Cu–BiOCl composite lm, which matches the position of the (111) plane of the facecentered cubic (fcc) Cu metal. Additional reections assigned to metallic Bi are also detected in both patterns. The detection of metallic Bi would be in agreement with (R4). The patterns were Rietveld tted and the obtained volume phase percentages, crystallite sizes and microstrains are listed in Table 1. Remarkably, all phases exhibit sub-100 nm hDi values,
3.2
Fig. 5
Mechanical properties
The mechanical properties of the electrodeposited layers were assessed by nanoindentation and compared to those of commercial bulk BiOCl. Representative load–displacement
XRD pattern of the (a) Cu-free BiOCl film and (b) Cu–BiOCl composite film.
Table 1 Phase volume percentages, crystallite sizes (hDi) and microstrains (h31/2i2) determined from the Rietveld fitting of the full XRD patterns for Cu-free and Cu– BiOCl composite films
hDi/nm (5)
% by vol
h31/2i2
Film
BiOCl
Bi
Cu
BiOCl
Bi
Cu
BiOCl
Bi
Cu
BiOCl Cu–BiOCl
59 43
41 22
— 35
83 65
28 20
— 105
8.2 104 2.0 103
2.0 104 5.0 104
— 8.0 104
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Fig. 7 Representative load (P)–displacement (h) nanoindentation curves of (a) BiOCl and (b) Cu–BiOCl composite films. A schematic picture of the indentation process is shown as an inset.
The porosity level is known to have a strong inuence on the elastic constants of metallic and ceramic materials. In a rst approximation, it has been shown that:41–43 rporous n Eporous ¼ (5) Ebulk rbulk where rporous/rbulk is the relative density. The exponent n is equal to 2 for open-celled foams or sponge-like structures,43,44 whereas n ¼ 3 for materials exhibiting an array of pores arranged forming a honeycomb hexagonal array normal to the surface.45 Given the porous nature of the electrodeposited layers, n 2 can be assumed in our case. Moreover, since the Young's modulus of diamond is very high, the true Young's moduli of Cu–BiOCl and BiOCl are very similar to the measured Er values (3). The relative density is in turn related to the porosity volume fraction: rporous ¼1P rbulk
Fig. 6 (a) TEM and (b and c) HRTEM images of the Cu–BiOCl composite. The inset in (a) shows the FFT of the region shown in (b), for which the spot of the (110) plane of tetragonal BiOCl has been indicated. The inset in (c) shows the FFT of the area enclosed in the red square, for which the (012) plane of rhombohedral metallic Bi has been indicated.
nanoindentation curves of the Cu-free and Cu–BiOCl deposits are shown in Fig. 7. The maximum indentation displacement attained in the porous BiOCl (curve a) is clearly larger than for the Cu–BiOCl composite (curve b), indicating that the latter is mechanically harder. Remarkably, not only the hardness, but also the reduced Young's modulus, the ratio H/Er (which is related to the wear resistance), and the ratio We/Wtot (indicative of the elastic recovery) are larger for the Cu-containing lm (see Table 2), thus evidencing the key role of Cu in enhancing the mechanical integrity of these electrodeposited materials.
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(6)
In a rst approximation, assuming Er z E, eqn (5) can be used to roughly estimate the porosity level of the electrodeposited BiOCl, taking the measured mechanical properties of commercial BiOCl as those of the bulk BiOCl. Considering a bulk density for BiOCl of 7.78 g cm3,46 using n ¼ 2, and Ebulk ¼ 0.41 EBi + 0.59 EBiOCl (in agreement with the rule of mixtures for the E value of a composite material),47 the density of the porous BiOCl lm is calculated to be 0.84 g cm3 (i.e., porosity level P ¼ 89%). This high porosity level explains the drastic decrease in the Young's modulus of the electroplated BiOCl lm with respect to the commercial material. If eqn (5) is used for Cu–BiOCl (and the bulk Young's modulus of Cu–BiOCl is assumed to be 0.35 ECu + 0.22 EBi + 0.43 EBiOCl) where EBiOCl is that of the commercial, then the porosity level for the Cu–BiOCl lm electrodeposited with the Triton X-100 concentration of 4 104 M would be P ¼ 0.87, which is similar to the value obtained in other porous electrodeposited systems, like localized electrodeposited Cu–Ni porous microcolumns.44 The
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Table 2 BiOCl
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Summary of the mechanical properties, evaluated by nanoindentation, of the electrodeposited Cu–BiOCl and BiOCl films, as well as commercial non-porous
Sample
H (GPa)
Er (GPa)
H/Er
We/Wtot
Cu–BiOCl BiOCl Commercial BiOCl (non-porous)
0.072 0.004 0.020 0.001 2.90 0.05
1.18 0.09 0.42 0.06 39.0 0.5
0.061 0.005 0.047 0.007 0.074 0.002
0.598 0.068 0.179 0.018 0.797 0.050
same method can be used to determine the porosity level of Cu– BiOCl lms electrodeposited using different Triton X-100 concentrations from the values of reduced Young's modulus (see Table S1†). The porosity level tends to increase as the Triton X-100 concentration decreases (i.e., P ¼ 88% for a Triton X-100 concentration of 2 104 M and P ¼ 92% for a Triton X-100 concentration of 1 104 M). It should be pointed out that this method of indirect assessment of the porosity degree from mechanical measurements has been corroborated in other materials by means of complementary techniques, such as Brunauer–Emmett–Teller (BET) analysis.48 Concerning the hardness, the value obtained in the porous BiOCl (H ¼ 20 MPa) is also signicantly lower than the hardness of commercial BiOCl (H ¼ 2.9 GPa, which is similar to the reported value for BiOCl from the literature).46 The decrease of hardness (or compressive yield stress) with porosity is also a well-documented effect42,49,50 and has been modeled using nite element simulations of nanoindentation curves.50 An equation analogous to eqn (5) may be used to correlate the yield stress of the porous structure with that of the bulk solid material: rporous m sporous ¼ C2 sbulk (7) rbulk
Fig. 8 3D CSLM images of the Cu–BiOCl composite film in (a) reflection mode and (b) photoluminescence response.
where C2 ¼ 0.3 and m ¼ 1.5 for an open-cell foam.42 Although the relationship between hardness and yield stress in bulk oxide ceramic is oen taken as Hbulk z 1.6sbulk,51 a quantitative relationship between hardness and yield stress in porous materials is yet to be established. Some authors consider that the indenter in porous structures is not constrained by the surrounding material regardless of eventual densication. Then, the nanoindentation would be equivalent to a uniaxial compression test and Hporous ¼ sporous.42 In our case and using the aforementioned assumptions, eqn (7) would give m z 1.27 for porous BiOCl, which is a quite reasonable value. Finally, the elastic recovery for porous Cu–BiOCl, Wel/Wtot ¼ 0.598, is not far from that of commercial BiOCl. This parameter indicates how much energy is released from the material aer being loaded and it could be of particular interest in applications subject to impact loading.52
3.3
Optical emission properties/photoluminescence
Three-dimensional photoluminescence (PL) and structure (reection mode) images of the porous Cu–BiOCl composite lm are shown in Fig. 8 (see also ESI, video le†). Remarkably, the PL response consists of broad-band emission which covers the entire visible range (in the wavelength interval 450–750 nm, as shown in Fig. 9) and has a maximum at around 503 nm (i.e.,
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Fig. 9 Spectral profiles representing the mean fluorescence intensity versus emission wavelength for Cu–BiOCl films in the 450 nm to 750 nm range. Data are the mean from n ¼ 10 regions of interest from three fields examined.
green region). The observed distribution of PL is homogeneous in depth. Unfortunately, the PL intensity of the Cu-free BiOCl lms was much lower, probably due to the presence of large amounts of Bi, which acts absorbing part of the incident light,
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Paper thus attenuating the PL response. The broad green emission observed in Cu–BiOCl is in agreement with previous studies on photoluminescence of ower-like BiOCl structures.53 Such green emission is shied to lower energies (around 2.5 eV), as compared to the energy bandgap of bulk BiOCl (3.46 eV).54 Therefore, the observed PL signal is mainly attributed to deep trapped uorescence emission arising from surface states, i.e. light emitted when a photogenerated hole, trapped in a midgap state, encounters an electron before it can relax non-radiatively to the ground state.55 As in many other photoluminescent materials there is a large Stokes shi between the absorbance and emission peaks, i.e., the emitted light is of lower energy than the excitation light. Such a difference stems from the nonradiative relaxation being lost as thermal decay. Also worth mentioning is that contrary to organic uorescent dots, where a pronounced decrease in the PL intensity is observed during illumination with UV or visible light (a phenomenon referred to as photobleaching56), here only a decrease of 20% of PL intensity was observed aer continuous irradiation for 10 min. Due to this stability in the PL signal and the relatively large width of the PL peak, the investigated material could be interesting for a variety of opto-electronic devices.
4.
Conclusions
3D hierarchically porous crack-free Cu–BiOCl composite lms have been fabricated by one-step electrodeposition in acidic electrolytes containing Triton X-100. The lms exhibit spherical macropores surrounded by nanoporous walls in which the BiOCl nanoplates can be clearly identied. It has been demonstrated that the electrodeposited Cu particles act as building blocks of the porous structure since less porous lms are obtained from Cu(II)-free solutions. The nanomechanical properties of the BiOCl and Cu–BiOCl lms correlate with the porous character of the lms. 3D CSLM images demonstrate a homogeneous distribution and stable photoluminescence within the 3D structure arising from the BiOCl compound. It is envisaged that this synthetic approach could be easily scaled to different combinations of metals and oxyhalide compounds to render 3D hierarchically porous structures for a variety of applications ranging from photonics, photocatalysis, sensing or electrocatalysis.
Acknowledgements Financial support from the MAT2011-27380-C02-01 research project from the Spanish MINECO and the 2009SGR-1292 from the Generalitat de Catalunya is acknowledged. The Servei d0 An` alisi Qu´ımica and the Servei de Difracci´o de la UAB are acknowledged. M.D.B. acknowledges funding support from an ICREA-Academia award.
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This journal is ª The Royal Society of Chemistry 2013
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PAPER
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Cite this: Nanoscale, 2013, 5, 12542
3D hierarchically porous Cu–BiOCl nanocomposite films: one-step electrochemical synthesis, structural characterization and nanomechanical and photoluminescent properties† ´,b Bradley J. Nelson,b Maria Dolors Baro ´, a Miguel Guerrero,*a Salvador Pane c d a ` nica Rolda ´n, Jordi Sort* and Eva Pellicer* Mo Three-dimensional (3D) hierarchically porous composite Cu–BiOCl films have been prepared by a facile onestep galvanostatic electrodeposition process from acidic electrolytic solutions containing Cu(II) and Bi(III) chloride salts and Triton X-100. The films show spherical, micron-sized pores that spread over the whole film thickness. In turn, the pore walls are made of randomly packed BiOCl nanoplates that are assembled leaving micro–nanopore voids beneath. It is believed that Cu grows within the interstitial spaces between the hydrogen bubbles produced from the reduction of H+ ions. Then, the BiOCl sheets accommodate in the porous network defined by the Cu building blocks. The presence of Cu tends to enhance the mechanical stability of the composite material. The resulting porous Cu–BiOCl films exhibit
Received 8th July 2013 Accepted 8th October 2013
homogeneous and stable-in-time photoluminescent response arising from the BiOCl component that spreads over the entire 3D porous structure, as demonstrated by confocal scanning laser microscopy. A broad-band emission covering the entire visible range, in the wavelength interval 450–750 nm, is
DOI: 10.1039/c3nr03491g
obtained. The present work paves the way for the facile and controlled preparation of a new
www.rsc.org/nanoscale
generation of photoluminescent membranes.
1.
Introduction
Three-dimensional (3D) porous lms have become increasingly popular in the last few years as they offer a number of applications in electrocatalysis,1 separation and concentration of gas molecules,2 batteries,3 sensors4 and electronic devices,5 among others. Several strategies are currently being pursued for the preparation of porous lms like de-alloying,6 galvanic replacement,7 or templating (e.g. using colloidal crystals as templates).8 Electrodeposition has recently emerged as a powerful and cheap technique to prepare 3D porous lms of metals and metallic alloys through different strategies: deposition in rigid porous templates (e.g. polystyrene opal templates),9 using hydrogen bubbles as a dynamic template10 or from electrolyte solutions containing suitable surfactants.11 The utilization of
Departament de F´ısica, Facultat de Ci`encies, Universitat Aut`onoma de Barcelona, E-08193 Bellaterra, Spain. E-mail: [email protected]; [email protected]
a
b
Institute of Robotics and Intelligent Systems (IRIS), Swiss Federal Institute of Technology (ETH) Zurich, CH-8092 Zurich, Switzerland
c
Servei de Microsc` opia, Universitat Aut`onoma de Barcelona, E-08193 Bellaterra, Spain
Instituci´ o Catalana de Recerca i Estudis Avançats (ICREA), Departament de F´ısica, Universitat Aut` onoma de Barcelona, E-08193 Bellaterra, Spain. E-mail: Jordi.Sort@ uab.cat
d
† Electronic supplementary 10.1039/c3nr03491g
information
12542 | Nanoscale, 2013, 5, 12542–12550
(ESI)
available.
See
DOI:
hydrogen co-evolution as a source of porosity during electrodeposition is a very convenient and fast way to prepare network porous structures of a number of metals (e.g. Cu,12 Ni,13 Pd,14 Pt,15 bimetallic Cu/Pd16) and alloys (e.g. Cu6Sn5).17 Among them, Cu is perhaps one of the most studied. Different electrolytes have been tested for the preparation of porous Cu lms with distinct pore morphologies, pore sizes and interconnectivities. Macroporous Cu lms constructed by nanodendritic walls have been obtained from acid sulphate solutions containing cetyltrimethylammonium bromide (CTAB).12 Cu foams with 3D interconnected spherical pore networks have been prepared in acid sulphate solutions containing HCl, polyethylene glycol (PEG) and 3-mercapto-1-propane sulfonic acid (MPSA) as additives.18 Remarkably, very few studies have reported the utilization of the hydrogen bubble template approach for the preparation of porous composite lms. A smart combination of metals with other compounds in a single 3D porous network structure could signicantly broaden the range of applications of porous lms. Bismuth oxychloride (BiOCl), a V–VI–VII group wide band gap semiconductor, is considered one of the most important bismuth oxyhalides. BiOCl has attracted an enormous amount of interest within the scientic community because of its excellent catalytic, optical, electric, magnetic and luminescent properties. As a result, BiOCl has found numerous uses as a
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Paper photocatalyst,19–21 or ferroelectric material,22,23 for many years. Moreover, it is widely used as a pigment in cosmetics, buttons and jewellery because it confers a semi-transparent appearance to the objects. Until now, several strategies have been reported for the preparation of BiOCl micro/nanostructures mostly in the form of powders: electrospinning,24 hydrothermal approach,25 reverse microemulsions26 or sonochemistry.27 However, for many industrial applications immobilized BiOCl lms are more adequate than powders. For example, problems encountered in nanopowder photocatalysts such as catalyst recovering (once the solution has been decontaminated) and particle aggregation can be circumvented if the active material can be properly immobilized without compromising its specic surface area. Nevertheless, studies on the preparation of nanostructured BiOCl lms are still scarce compared to the vast literature on BiOCl nanopowders. Cao and co-workers fabricated a BiOCl lm with ower-like hierarchical structure by reacting a Bi lm with H2O2/HCl solution.28 Very recently the preparation of BiOCl lms by electrodeposition of Bi and subsequent anodic oxidation has been reported.29 In all these studies the BiOCl lm is used for photocatalytic purposes. Other applications would also benet from having immobilized BiOCl lms like electrochemical sensing or optics. Herein, we present a facile, efficient and environmentally friendly approach for fabricating three-dimensional (3D) porous composite lms made of Cu and BiOCl components in one-step cathodic electrodeposition. An optimized bath formulation containing suitable concentrations of bismuth(III) and copper(II) chloride salts and Triton X-100 as an additive yields a crack-free porous composite structure. Also, the mechanical properties of the composite lms have been studied by nanoindentation in order to correlate the obtained data with both the porosity degree and the presence of Cu in the lms. Finally, the hierarchically porous Cu–BiOCl lms exhibit a strong and stable photoluminescence response in the entire visible range.
2.
Experimental section
2.1
Preparation of Cu-free and Cu–BiOCl composite lms
Cu-free and Cu–BiOCl composites were prepared by direct current electrodeposition (ED) in a thermostatized onecompartment three-electrode cell using a PGSTAT302N Autolab potentiostat/galvanostat (Ecochemie). The working electrode was positioned vertically within the electrolyte and consisted of 5 mm 6 mm Si chips (crystal orientation 100), on top of which a Ti adhesion layer of 50 nm and a Au seed-layer of 500 nm had been successively deposited through e-beam evaporation. The working area was 0.25 cm2. A double junction AgrAgCl (E ¼ +0.210 V/SHE) reference electrode (Metrohm AG) was used with 3 M potassium chloride (KCl) inner solution and an interchangeable outer solution. The outer solution was made of 1 M sodium chloride. A platinum sheet served as the counter electrode. For Cu–BiOCl composite electrodeposition, the electrolyte contained 1.2 M HCl, 8 103 M BiCl3, 8 103 M CuCl2, 2 102 M NaCl, and variable amounts of Triton X-100, saccharine and CTAB. The Cu-free BiOCl lms were prepared This journal is ª The Royal Society of Chemistry 2013
Nanoscale from an electrolyte free of CuCl2 salt. Analytical grade reagents and ultrapure water (18 MU cm) were used to prepare the electrolyte. The cell was lled with 50 mL electrolyte. All the samples were obtained galvanostatically by applying a constant current density of 1 A cm2. Deposition was conducted at room temperature under vigorous stirring (5000 rpm) using a magnetic stirrer bar. A similar electrical charge owed in all experiments, yielding lm thicknesses between 75 and 100 mm. Prior to deposition, the Au surface was degreased with acetone and dried with a so tissue. The Si backside was insulated by painting it with a nonconductive ink to ensure only the Au surface was conductive. Aer deposition, the coated substrates were rst dipped in 0.12 M HCl, aerwards dipped in Milli-Q water, and nally dried in air. 2.2
Morphological and structural characterization
The morphology and structure of the deposits was studied by scanning electron microscopy (SEM), high resolution transmission electron microscopy (HRTEM) and X-ray diffraction (XRD). SEM imaging was performed on a Zeiss Merlin microscope. TEM characterisation was carried out on a JEOL JEM2011 microscope operated at 200 kV. For TEM observations, a small area of sample surface was scratched-off and the powder released was dispersed in ethanol. Aerwards, one or two drops of the suspension were placed dropwise onto a holey carboncoated Cu TEM grid. The chemical composition of the lms was determined by energy dispersive X-ray spectroscopy (EDX). The pore size distributions were determined by manual analysis of enlarged micrographs by measuring ca. 100 pores on a given grid to obtain a statistical size distribution and a mean pore diameter. XRD patterns were recorded on a Philips X'Pert diffractometer using Cu Ka radiation, in the 20–90 2q range (0.03 step size, 8 s holding time). The global structural parameters, such as phase percentages, crystallite sizes, hDi (dened here as the average coherently diffracting domain sizes), and microstrains or atomic level deformations, h32i1/2, were evaluated by tting the full XRD patterns using the “Materials Analysis Using Diffraction” (MAUD) Rietveld renement program.30–32 Fourier transform infrared spectroscopy (FT-IR) was carried out on a Bruker TENSOR27 instrument. 2.3
Mechanical characterization
The mechanical properties (hardness, reduced elastic modulus and elastic recovery) were evaluated by means of nanoindentation, using a UMIS device from Fischer-Cripps Laboratories equipped with a Berkovich pyramidal-shaped diamond tip, operating in the load control mode. The maximum value of applied load was 0.6 mN to ensure that the lateral size of the indentation imprint was much smaller than the lm thickness. The thermal dri during nanoindentation was lower than 0.05 nm s1. Proper corrections for the contact area (calibrated with a fused quartz specimen), instrument compliance, and initial penetration depth (a pre-contact load of 0.2 mN was used) were applied. The hardness (H) and reduced elastic modulus (Er) values were derived from the load–displacement
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curves using the method of Oliver and Pharr.33 From the initial unloading slope, the contact stiffness, S, was determined as: S¼
dP dh
(1)
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where P and h denote the applied force and the penetration depth during nanoindentation, respectively. The reduced Young's modulus was evaluated based on its relationship with the contact area, A, and contact stiffness: pffiffiffiffi 2 S ¼ b pffiffiffiffi Er A p
(2)
where b is King's factor that depends on the geometry of the indenter (b ¼ 1.034 for a Berkovich indenter).34 The reduced modulus takes into account the elastic displacements that occur in both the specimen, with Young's modulus E and Poisson's ratio n, and the diamond indenter, with elastic constants Ei and ni (for diamond, Ei ¼ 1140 GPa and ni ¼ 0.07): 1 1 n2 1 ni 2 þ ¼ Er E Ei
(3)
The hardness was calculated from the following expression: PMax H¼ A
(4)
where PMax is the maximum load applied during nanoindentation. Finally, the elastic recovery was evaluated as the ratio between the elastic and total (plastic + elastic) energies during nanoindentation, Wel/Wtot. The values of Wel were calculated as the area between the unloading curve and the displacement axis. In turn, Wtot is the area between the loading curve and the displacement axis.34,35 To extract semi-quantitative information on the porosity level of the electrodeposited lms, the mechanical properties of non-porous bulk commercial BiOCl (Sigma-Aldrich, 98% purity) were also measured by nanoindentation. The results presented in this work are the statistical average of a set of 75 indentations for each sample. 2.4
Photoluminescence measurements
The photoluminescence (PL) properties of the lms were studied by Confocal Scanning Laser Microscopy (CSLM) analysis. Samples were mounted on Ibidi culture dishes (Ibidi GmbH, Martinsried, Germany) and imaged with a TCS-SP5 CSLM microscope (Leica Microsystems CMS GmbH Mannheim, Germany) using a Plan Apochromat 10/0.4 and 20/0.7 (dry) objectives. Both Cu-free and Cu–BiOCl composites were excited with an argon laser (488 nm) and the autouorescence was detected in the 500–785 nm range. The three-dimensional (3D) structure of lms was imaged using the same confocal microscope in reection mode with a 488 nm argon laser and detected in the 480–495 nm range. To determine the 3D microstructure, stacks of 50 to 80 sections were collected every 1 mm along the material's thickness. Three-dimensional models were generated from the xyz series using the Imaris X64 v. 6.2.0 soware (Bitplane; Z¨ urich, Switzerland). A series of images (xyl) were taken to determine the spectral emissions of the samples and to establish their maxima.
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Intrinsic uorescence was excited with a 405 nm line of a blue diode laser. Fluorescence emission was collected with a hybrid detector in 20 nm bandwidth increments (lambda step size ¼ 7 nm) in the range from 450 nm to 750 nm. A set of 30 regions of interest (ROIs) of 100 mm2 was used to analyze the mean uorescence intensity (MFI) of the samples in relation to the emission wavelength. Finally, photostability experiments were performed to monitor long time-lapse experiments. A 50 mW blue diode 405 nm laser was used at a fairly high output (Acousto-Optic Tunable Filter, AOTF ¼ 90%) in a eld of 1.5 mm 1.5 mm and the uorescence signal was detected in the range from 450 nm to 750 nm. Images were taken every 1 s for 10 min. A set of 20 ROIs of 100 mm2 were selected to show the MFI in the region in relation to time. Data from all studies were analyzed using the LAS AF soware ver. 2.4.1. (Leica Microsystems).
3.
Results and discussion
3.1 Synthesis, morphological and structural characterization of Cu-free and Cu–BiOCl composite lms Representative on-top SEM images of the Cu–BiOCl composite lms grown in the electrolyte containing 4 104 M of Triton X-100 are shown in Fig. 1. At low magnication, a uniform distribution of spherical pores of 5–10 mm in diameter (macroporosity) is observed over the entire surface of the lm (Fig. 1a and b). Moreover, the pore walls are not dense but inherently porous, with voids having dimensions from tenths of nanometers to a few microns (Fig. 1c). A close-up observation of the pore walls (Fig. 1d) reveals that they are in fact built from numerous randomly packed nanoplates, and that the nanoplates are stacked together either through plane-to-plane or edge-to-edge conjunctions, giving a ower-like appearance. Hence, the lms are hierarchically porous. The porosity spreads over the whole thickness as demonstrated by the cross-section image of the composite layer (Fig. 2). It is worth mentioning that the Triton X-100 additive suppresses the formation of cracks. With other additives like saccharine and CTAB cracked porous lms were obtained (ESI, Fig. S1†). As a result, the samples were so powdery and brittle that good contact between
Fig. 1 On-top SEM images of the Cu–BiOCl composite obtained in the electrolyte with 4 104 M Triton X-100 at different magnifications.
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Nanoscale
Fig. 2 Cross-section SEM image of the Cu–BiOCl composite obtained in the electrolyte with 4 104 M Triton X-100.
the samples and the indenter tip could not be attained and the nanoindentation data became noisy and unreliable. For this reason, samples prepared using saccharine or CTAB as additives were not considered for in-depth mechanical and photoluminescence studies. Saccharine is typically used as a stressreliever and grain-rener in electrodeposited metals and alloys,35 but does not work for all systems and it has little effect on porosity. CTAB is a cationic surfactant with a positively charged head group that interacts electrostatically with the electrode surface. It gets adsorbed on the electroactive sites and, as a consequence, changes crystal growth mode and, in turn, grain morphology and size.36,37 Triton X-100 is recognized as an excellent wetting agent, which means that it favours H2 bubble detachment from the substrate,38 which likely promotes the evolution of porosity in the coating. At the same time, it would inhibit the formation of cracks in the lm as it builds. It is believed that such a synergetic effect is not achieved with saccharine and CTAB additives. Signicantly, Cu-free BiOCl lms exhibit different pore architecture (Fig. 3a). The lm surface is still very rough but the spherical macropores are no longer evident (Fig. 3b). A rose-like morphology is observed at high magnication (Fig. 3c). This suggests that the reduction of H+ (R1) that accompanies Cu discharge (R2) is responsible for the formation of the spherical macropores. The BiOCl nanoplates are then accommodated within the macropore walls dened by the Cu building blocks. The electrochemical reactions involved in the formation of the macropore skeleton are thus: H+ + 1e / ½H2
(R1)
Cu2+ + 2e / Cu
(R2)
On the other hand, the formation of BiOCl nanoplates would mainly proceed via the reaction of Bi(III) with an electrogenerated base (OH) in the presence of Cl (R3): Bi3+ + Cl + 2OH / BiOCl + H2O
(R3)
The formation of the BiOCl component does not take place by simply dipping the substrate in the electrolyte if no current is owing through the system (i.e., by hydrolysis). Hence, the formation of the BiOCl nanoplates necessarily involves a reaction with any species generated during the electrolysis (could be
This journal is ª The Royal Society of Chemistry 2013
Fig. 3
(a–c) On-top SEM images of the Cu-free BiOCl film.
H2O2 that reacts with freshly reduced Bi, as well). Moreover, the electroreduction of Bi(III) toward metallic Bi would also take place: Bi3+ + 3e / Bi
(R4)
It has been reported that metallic bismuth electrodeposited at hydrogen evolution potentials exhibits a non-hierarchical porous dendritic structure.39 The potential–time (E–t) transients recorded during Cu-free and Cu–BiOCl lms are shown in the ESI, Fig. S2.† The E–t curves show an important decrease of the deposition potential toward more positive values during the rst few seconds (nucleation events) until stabilization around 3 V vs. Ag/AgCl is achieved. The stabilized potential is similar for both Cu-free and Cu–BiOCl lms and there is not any signicant decrease/increase of the average potential with time, which suggests that the surface area of the deposition front remains almost constant. The uctuations are due to the intense hydrogen evolution. Both the macropore size and size distribution of the Cu–BiOCl hierarchically layered composites can be greatly tuned by varying the Triton X-100 concentration in the electrolyte (ESI, Fig. S3†). Namely, both values decrease as the concentration of Triton X-100 is increased from 1 104 M (mean macropore size 24.7 mm) to 4 104 M (mean macropore size 9.1 mm). Energy-dispersive X-ray spectroscopy (EDX) analysis shows the presence of Bi, O and Cl (coming from the BiOCl layer) as well as Cu homogeneously distributed within the nanocomposite (Fig. 4).
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Fig. 4
Paper
EDX mapping distribution of Cu, Bi, O and Cl elements in the Cu–BiOCl composite film.
evidencing the nanocrystalline character of the composite electrodeposited lms. In turn, the values of microstrains (which are related to the density of dislocations and other structural defects) are relatively low, much lower than those typically obtained in heavily deformed metals,40 but still higher than the microstrain values in some electrodeposited nanocrystalline metallic alloys.35 To gain further information on Cu–BiOCl hierarchically porous lms, the product was further investigated by HRTEM. Fig. 6a shows a blend of nanosheets corresponding to BiOCl. The size of the sheets is well in the sub-100 nm range, in agreement with XRD results. Fig. 6b is a lattice-resolved HRTEM image taken on a single nanosheet. The corresponding lattice fringes have an interplanar distance of 0.27 nm, which matches the d-spacing of the (110) reection of tetragonal BiOCl (d110 ¼ 0.2753 nm). This is in accordance with the selected-area electron diffraction (SAED), indexed to a pure tetragonal phase (inset of Fig. 6a). This result indicates that the preponderant growth direction of the BiOCl layer is the [001] orientation, which is parallel to the (110) orientation. Metallic Bi on the BiOCl nanosheets is also detected, in agreement with XRD data (Fig. 6c). The inset shows the FFT of the area enclosed in red that belongs to the rhombohedral Bi phase (JCPDS no. 44-1246).39 The IR spectra of the sample show the typical bands of BiOCl in three different regions (2854 cm1, 1463 cm1 and 612 cm1) (ESI, Fig. S4†), again conrming the presence of the BiOCl component.
The crystallographic structure of the as-prepared Cu-free and Cu–BiOCl coatings was studied using XRD (Fig. 5). Apart from the reections coming out from the substrates (Si and Au), some of the diffraction peaks can be assigned to the tetragonal BiOCl phase (S.G., P4/nmm; JCPDS no. 06-0249, a ¼ b ¼ 0.3891 nm and c ¼ 0.7369 nm). In addition, a narrower reection at 2q ¼ 43.30 is identied in the XRD pattern of the Cu–BiOCl composite lm, which matches the position of the (111) plane of the facecentered cubic (fcc) Cu metal. Additional reections assigned to metallic Bi are also detected in both patterns. The detection of metallic Bi would be in agreement with (R4). The patterns were Rietveld tted and the obtained volume phase percentages, crystallite sizes and microstrains are listed in Table 1. Remarkably, all phases exhibit sub-100 nm hDi values,
3.2
Fig. 5
Mechanical properties
The mechanical properties of the electrodeposited layers were assessed by nanoindentation and compared to those of commercial bulk BiOCl. Representative load–displacement
XRD pattern of the (a) Cu-free BiOCl film and (b) Cu–BiOCl composite film.
Table 1 Phase volume percentages, crystallite sizes (hDi) and microstrains (h31/2i2) determined from the Rietveld fitting of the full XRD patterns for Cu-free and Cu– BiOCl composite films
hDi/nm (5)
% by vol
h31/2i2
Film
BiOCl
Bi
Cu
BiOCl
Bi
Cu
BiOCl
Bi
Cu
BiOCl Cu–BiOCl
59 43
41 22
— 35
83 65
28 20
— 105
8.2 104 2.0 103
2.0 104 5.0 104
— 8.0 104
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Paper
Nanoscale
Fig. 7 Representative load (P)–displacement (h) nanoindentation curves of (a) BiOCl and (b) Cu–BiOCl composite films. A schematic picture of the indentation process is shown as an inset.
The porosity level is known to have a strong inuence on the elastic constants of metallic and ceramic materials. In a rst approximation, it has been shown that:41–43 rporous n Eporous ¼ (5) Ebulk rbulk where rporous/rbulk is the relative density. The exponent n is equal to 2 for open-celled foams or sponge-like structures,43,44 whereas n ¼ 3 for materials exhibiting an array of pores arranged forming a honeycomb hexagonal array normal to the surface.45 Given the porous nature of the electrodeposited layers, n 2 can be assumed in our case. Moreover, since the Young's modulus of diamond is very high, the true Young's moduli of Cu–BiOCl and BiOCl are very similar to the measured Er values (3). The relative density is in turn related to the porosity volume fraction: rporous ¼1P rbulk
Fig. 6 (a) TEM and (b and c) HRTEM images of the Cu–BiOCl composite. The inset in (a) shows the FFT of the region shown in (b), for which the spot of the (110) plane of tetragonal BiOCl has been indicated. The inset in (c) shows the FFT of the area enclosed in the red square, for which the (012) plane of rhombohedral metallic Bi has been indicated.
nanoindentation curves of the Cu-free and Cu–BiOCl deposits are shown in Fig. 7. The maximum indentation displacement attained in the porous BiOCl (curve a) is clearly larger than for the Cu–BiOCl composite (curve b), indicating that the latter is mechanically harder. Remarkably, not only the hardness, but also the reduced Young's modulus, the ratio H/Er (which is related to the wear resistance), and the ratio We/Wtot (indicative of the elastic recovery) are larger for the Cu-containing lm (see Table 2), thus evidencing the key role of Cu in enhancing the mechanical integrity of these electrodeposited materials.
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(6)
In a rst approximation, assuming Er z E, eqn (5) can be used to roughly estimate the porosity level of the electrodeposited BiOCl, taking the measured mechanical properties of commercial BiOCl as those of the bulk BiOCl. Considering a bulk density for BiOCl of 7.78 g cm3,46 using n ¼ 2, and Ebulk ¼ 0.41 EBi + 0.59 EBiOCl (in agreement with the rule of mixtures for the E value of a composite material),47 the density of the porous BiOCl lm is calculated to be 0.84 g cm3 (i.e., porosity level P ¼ 89%). This high porosity level explains the drastic decrease in the Young's modulus of the electroplated BiOCl lm with respect to the commercial material. If eqn (5) is used for Cu–BiOCl (and the bulk Young's modulus of Cu–BiOCl is assumed to be 0.35 ECu + 0.22 EBi + 0.43 EBiOCl) where EBiOCl is that of the commercial, then the porosity level for the Cu–BiOCl lm electrodeposited with the Triton X-100 concentration of 4 104 M would be P ¼ 0.87, which is similar to the value obtained in other porous electrodeposited systems, like localized electrodeposited Cu–Ni porous microcolumns.44 The
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Table 2 BiOCl
Paper
Summary of the mechanical properties, evaluated by nanoindentation, of the electrodeposited Cu–BiOCl and BiOCl films, as well as commercial non-porous
Sample
H (GPa)
Er (GPa)
H/Er
We/Wtot
Cu–BiOCl BiOCl Commercial BiOCl (non-porous)
0.072 0.004 0.020 0.001 2.90 0.05
1.18 0.09 0.42 0.06 39.0 0.5
0.061 0.005 0.047 0.007 0.074 0.002
0.598 0.068 0.179 0.018 0.797 0.050
same method can be used to determine the porosity level of Cu– BiOCl lms electrodeposited using different Triton X-100 concentrations from the values of reduced Young's modulus (see Table S1†). The porosity level tends to increase as the Triton X-100 concentration decreases (i.e., P ¼ 88% for a Triton X-100 concentration of 2 104 M and P ¼ 92% for a Triton X-100 concentration of 1 104 M). It should be pointed out that this method of indirect assessment of the porosity degree from mechanical measurements has been corroborated in other materials by means of complementary techniques, such as Brunauer–Emmett–Teller (BET) analysis.48 Concerning the hardness, the value obtained in the porous BiOCl (H ¼ 20 MPa) is also signicantly lower than the hardness of commercial BiOCl (H ¼ 2.9 GPa, which is similar to the reported value for BiOCl from the literature).46 The decrease of hardness (or compressive yield stress) with porosity is also a well-documented effect42,49,50 and has been modeled using nite element simulations of nanoindentation curves.50 An equation analogous to eqn (5) may be used to correlate the yield stress of the porous structure with that of the bulk solid material: rporous m sporous ¼ C2 sbulk (7) rbulk
Fig. 8 3D CSLM images of the Cu–BiOCl composite film in (a) reflection mode and (b) photoluminescence response.
where C2 ¼ 0.3 and m ¼ 1.5 for an open-cell foam.42 Although the relationship between hardness and yield stress in bulk oxide ceramic is oen taken as Hbulk z 1.6sbulk,51 a quantitative relationship between hardness and yield stress in porous materials is yet to be established. Some authors consider that the indenter in porous structures is not constrained by the surrounding material regardless of eventual densication. Then, the nanoindentation would be equivalent to a uniaxial compression test and Hporous ¼ sporous.42 In our case and using the aforementioned assumptions, eqn (7) would give m z 1.27 for porous BiOCl, which is a quite reasonable value. Finally, the elastic recovery for porous Cu–BiOCl, Wel/Wtot ¼ 0.598, is not far from that of commercial BiOCl. This parameter indicates how much energy is released from the material aer being loaded and it could be of particular interest in applications subject to impact loading.52
3.3
Optical emission properties/photoluminescence
Three-dimensional photoluminescence (PL) and structure (reection mode) images of the porous Cu–BiOCl composite lm are shown in Fig. 8 (see also ESI, video le†). Remarkably, the PL response consists of broad-band emission which covers the entire visible range (in the wavelength interval 450–750 nm, as shown in Fig. 9) and has a maximum at around 503 nm (i.e.,
12548 | Nanoscale, 2013, 5, 12542–12550
Fig. 9 Spectral profiles representing the mean fluorescence intensity versus emission wavelength for Cu–BiOCl films in the 450 nm to 750 nm range. Data are the mean from n ¼ 10 regions of interest from three fields examined.
green region). The observed distribution of PL is homogeneous in depth. Unfortunately, the PL intensity of the Cu-free BiOCl lms was much lower, probably due to the presence of large amounts of Bi, which acts absorbing part of the incident light,
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Paper thus attenuating the PL response. The broad green emission observed in Cu–BiOCl is in agreement with previous studies on photoluminescence of ower-like BiOCl structures.53 Such green emission is shied to lower energies (around 2.5 eV), as compared to the energy bandgap of bulk BiOCl (3.46 eV).54 Therefore, the observed PL signal is mainly attributed to deep trapped uorescence emission arising from surface states, i.e. light emitted when a photogenerated hole, trapped in a midgap state, encounters an electron before it can relax non-radiatively to the ground state.55 As in many other photoluminescent materials there is a large Stokes shi between the absorbance and emission peaks, i.e., the emitted light is of lower energy than the excitation light. Such a difference stems from the nonradiative relaxation being lost as thermal decay. Also worth mentioning is that contrary to organic uorescent dots, where a pronounced decrease in the PL intensity is observed during illumination with UV or visible light (a phenomenon referred to as photobleaching56), here only a decrease of 20% of PL intensity was observed aer continuous irradiation for 10 min. Due to this stability in the PL signal and the relatively large width of the PL peak, the investigated material could be interesting for a variety of opto-electronic devices.
4.
Conclusions
3D hierarchically porous crack-free Cu–BiOCl composite lms have been fabricated by one-step electrodeposition in acidic electrolytes containing Triton X-100. The lms exhibit spherical macropores surrounded by nanoporous walls in which the BiOCl nanoplates can be clearly identied. It has been demonstrated that the electrodeposited Cu particles act as building blocks of the porous structure since less porous lms are obtained from Cu(II)-free solutions. The nanomechanical properties of the BiOCl and Cu–BiOCl lms correlate with the porous character of the lms. 3D CSLM images demonstrate a homogeneous distribution and stable photoluminescence within the 3D structure arising from the BiOCl compound. It is envisaged that this synthetic approach could be easily scaled to different combinations of metals and oxyhalide compounds to render 3D hierarchically porous structures for a variety of applications ranging from photonics, photocatalysis, sensing or electrocatalysis.
Acknowledgements Financial support from the MAT2011-27380-C02-01 research project from the Spanish MINECO and the 2009SGR-1292 from the Generalitat de Catalunya is acknowledged. The Servei d0 An` alisi Qu´ımica and the Servei de Difracci´o de la UAB are acknowledged. M.D.B. acknowledges funding support from an ICREA-Academia award.
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