Preparing the Way for Doping Wurtzite Silicon Nanowires while

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Preparing the Way for Doping Wurtzite Silicon Nanowires while Retaining the Phase Filippo Fabbri,†,‡ Enzo Rotunno,† Laura Lazzarini,† Daniela Cavalcoli,§ Antonio Castaldini,§ Naoki Fukata,‡ Keisuke Sato,§ Giancarlo Salviati,*,† and Anna Cavallini§ †

IMEM-CNR Institute, University Campus, Viale Delle Scienze 37/A, I-43124 Parma, Italy International Center for Materials Nanoarchitectonics, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan § Department of Physics and Astronomy, University of Bologna, Viale Berti Pichat 6/2, 40127 Bologna, Italy ‡

S Supporting Information *

ABSTRACT: It is demonstrated that boron-doped nanowires have predominantly long-term stable wurtzite phase while the majority of phosphorus-doped ones present diamond phase. A simplified model based on the different solubility of boron and phosphorus in gold is proposed to explain their diverse effectiveness in retaining the wurtzite phase. The wurtzite nanowires present a direct transition at the Γ point at approximately 1.5 eV while the diamond ones have a predominant emission around 1.1 eV. The aforementioned results are intriguing for innovative solar cell devices. KEYWORDS: Wurtzite nanowires, silicon, HRTEM, cathodoluminescence, SPV

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properties and bandgap of wurtzite gold catalyzed SiNW mats, are strongly affected by B- (p-type) and P- (n-type) doping. In particular, we report on the solid experimental evidence that B and P dopants have different effects on retaining the phase of the grown NWs. The great majority of undoped NWs shows totally wurtzite phase, which is still present after three years time. B-doped NWs in the mats predominantly keep the wurtzite phase meanwhile P-doped ones have mainly diamond phase. The influence of the wurtzite phase on the optical and optoelectronic properties of the NWs is investigated by surface photovoltage spectroscopy (SPS),13,14 cathodoluminescence spectroscopy (CL), 15 and spectral photoconductivity (PC)14,16 techniques, which are very efficient tools to study near-band-edge (NBE) absorption and emission properties. We demonstrate that B-doped wurtzite NWs possess an optical band gap at around 1.5 eV, wider than diamond silicon, while the P-doped diamond ones present the bandgap value of diamond bulk Si. We refer to bulk silicon because the minimum observed NW diameter is larger than the size value reported in the literature for quantum confinement effects in diamond silicon nanocrystals (see in the following). Besides, we show that the B-doped NWs provide a remarkably lower reflectivity in the visible/near-infrared region that can be related to a combined effect of morphological and structural properties. To

ne-dimensional semiconductor architectures, such as nanowires1−3 and nanotubes,4 are crucially important in materials-based applications requiring large surface area, morphological control, and superior charge transport. In particular, diamond silicon nanowires (Si NWs) have recently gained increasing attention for solar cell applications due to their superior properties, such as diameter-tunable optical absorption5 and optical bandgap,6 low reflectivity,5,7,8 and stronger optical absorption5,7,8 in the visible region compared to bulk and amorphous Si. In addition, as demonstrated by Garnett and Yang,9 the use of diamond Si NWs increases the light-trapping properties over the randomized scattering limit. Recently, some of the authors have successfully grown p- and ntype diamond Si NWs by impurity doping respectively boron (B) and phosphorus (P), by laser ablation of a Si target.10 Wurtzite Si NWs are more attractive for solar cell applications due to the possible optical bandgap engineering of the different Si polymorphs in the device design, as for III−V based multijunctions solar cells. In the design of solar cells based on wurtzite NWs, the doping is of crucial importance to enable the control of conductivity and the formation of p/n junctions. The study of wurtzite NWs is still in its infancy and until now only a very few papers reporting about growth mechanisms and structural properties can be found in the literature.11,12 In this work, we grow gold-catalyzed silicon NWs with an electrically active P concentration estimated to be about 1019 cm−3 while that of the B is around 1018−19 cm−3. We demonstrate that the crystalline structure, hence the optical © 2013 American Chemical Society

Received: July 30, 2013 Revised: October 28, 2013 Published: November 13, 2013 5900

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A commercial Gatan monoCL system, fitted onto an S360 Cambridge SEM, is used for the CL spectroscopy. The system is equipped by a grating and a Ge photodiode, sensitive in the range 700−1550 nm (1.7−0.8 eV). CL experiments are performed at room temperature on tens of NWs after removing from the substrate. An accelerating voltage of 7.5 keV is chosen to release homogeneous amount of energy along the NW diameter (see Figure S10 and S13 in the Supporting Information). The current is 115 nA. For the CL analysis of bulk Si, the accelerating voltage is kept constant and the current is increased up to 10 μA due to its extremely lower radiative efficiency. The reflectance is measured at a reflection angle of 45° at room temperature. PC measurements are carried out at room temperature by means of a custom-made apparatus based on a CornerStone 260 monochromator equipped with a QTH lamp. The photocurrent signal is collected through two ohmic contacts on the NW mat by a lock-in amplifier and normalized to the photon flux. The ohmic contacts are prepared by evaporating a 300 nm thick gold onto the NW mats. The schematic of the sample and contacts for PC measurements is reported in Figure S14 in the Supporting Information. It is worth noting that, apart for the HRTEM observations, the results obtained come from the majority of the NWs in the mats, representing the behavior of wurtzite NWs in B-doped samples and the behavior of the diamond NWs in P-doped samples. The structural characterization is obtained by comparing high-resolution transmission electron microscopy (HRTEM) and electron diffraction on hundreds of NWs for different wire orientation. The TEM investigations show that the undoped, the B- and P-doped NWs can crystallize in the wurtzite or diamond phases and the two different polymorphs never occur in the same NW. From the TEM analyses, we found that the doping species have a completely different effect on the crystalline properties of the NWs, thereby causing the variation of the relative abundance of two polymorphs. The onset of the small percentage of the diamond phase in B-doped NWs compared to the undoped ones is probably related to the stress release caused by the formation of planar defects, that is, twin planes (see Figure S2 in the Supporting Information). All the NWs studied in this work are single crystals either wurtzite or diamond in nature. In these monocrystalline NWs, we do not observe any other of the possible hexagonal polymorphs intermediate between the diamond and the wurtzite structures (see, for example, ref 23). Moreover, according to the classification of Fontcuberta i Morral et al.11 we mainly observe the type A wurtzite structure with a = 0.38 nm and c = 0.627 nm and only occasionally the type B one with a = 0. 404 nm and c = 0.660 nm. The WZ phase is also confirmed by X-ray diffraction studies performed on borondoped samples (see Supporting Information for a deep discussion of this argument and Figures S2−S5). The type A phase is precisely determined by high-resolution transmission electron microscopy and diffraction and the observations are supported by computer image simulations. The need of image simulations comes from a renowned problem, that is, the chance that the hexagonal diffraction pattern of the wurtzite lattice observed in the projection [0001] could be mistaken with the [111] projection of a cubic crystal that is twinned on a (111) composition plane. An example is given in Figure 2a where a micrograph of 10 nm thin wurtzite silicon nanowires is reported. The red squared region is

the best of our knowledge, this is the first experimental evidence of B and P doping of wurtzite silicon NWs. Our results prepare the way to the control of structural, optical, and transport properties of Si nanowires in view of innovative device engineering. Si NWs are synthesized by means of a gold nanoclustercatalyzed vapor−liquid−solid (VLS) growth mechanism17 in a commercial cold wall CVD system at a total pressure ranging from 8 to 10 Torr. The growth rate of NWs was about 300 nm/ min. The substrates were thermally oxidized Si with an oxide thickness of 200 nm; Au nanocolloidal particles of 3 nm in diameter were used as metal catalyst. For B- and P-doped NWs synthesis, they are directly grown onto thermally oxidized Si substrates at a temperature of 600 °C for 30 min in flowing 19 sccm of SiH4 as the silicon reactant gas, 1% B2H6 (0.5 sccm) in H2 as p-type dopant gas, 1% PH3 in H2 as n-type dopant gas and 30 sccm of nitrogen (N2) as the carrier gas. It is important to remember that the growth temperature of our NWs is higher than that reported, for example, in refs 18−21 (that was around 450−460 °C), which present cubic Si NWs. In addition, also our precursors flow rates were different. On the contrary, our growth conditions are similar to those of ref 11, which reports on hexagonal NWs. Figure 1 shows the typical nanowire mat

Figure 1. SEM micrograph of a typical mat of B-doped Si NWs.

obtained by this procedure. The acceptor concentration in Bdoped NWs is roughly estimated to be in the order of 1018− 1019 cm−3 by Raman spectroscopy.10 The electrical active P concentration can be estimated to be about 1019 cm−3. The crystalline structure of Si NWs is determined by electron microscopy techniques in a JEOL 2200FS field emission microscope operating at 200 kV. The samples for HRTEM observation are prepared by dropping dilute dispersion solution of NWs onto holey-carbon grids, and their analysis is carried out on several hundreds of NWs. The optical and optoelectronic properties of the NWs are investigated by SPS, PC, CL, and reflectivity analyses. SPS measurements are conducted at room temperature by means of a custom-made apparatus based on a SPEX 500 M monochromator equipped with quartz-tungsten-halogen (QTH) lamp.22 The surface photovoltage (SPV) signal, capacitively picked up by a semitransparent electrode, goes through a high impedance field effect transistor preamplifer, and is measured by a lock-in amplifer. The SPV data are normalized by the photon flux impinging over the sample. 5901

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Figure 2. (a) HRTEM image of a 10 nm thick B-doped WZ Si NW (defocus = −400 nm); (b) magnification of the squared region in (a); (c,d) image simulation of a cubic twinned Si and a WZ silicon structure.

Figure 3. (a) Abundance of the WZ and diamond Si polymorphs NW in the undoped (green), B-doped (red), and P-doped (blue) samples. (b) Zero-loss HRTEM micrograph of a WZ B-doped NW taken in the [0001] zone axis. The growth direction is reported and indexed according the FFT shown in the inset. (c) Zero-loss high-resolution TEM micrograph of a P-doped diamond NW in the [011] zone axis. The growth direction is reported and indexed according the diffraction pattern shown in the inset. Short staking faults are present, circled in white. (d) Micro Raman spectra of the B- and P-dopes NWs compared to bulk silicon. The asymmetric broadening toward the higher wavenumber side in the B-doped spectrum is ascribed to the Fano effect due to B doping.10,29

enlarged in Figure 2b and is compared with the computer image simulations24 of the wurtzite silicon and the cubic twinned silicon, respectively, shown in Figure 2c,d, obtained in the same experimental conditions of Figure 2b. It is apparent that there is no possibility of mistaking the two structures. Our experimental image is exactly reproduced by the WZ structure simulation with lattice periodicity d10−10 = 0.33 nm. The same periodicity is also present in the image simulation of the twinned crystal but only as an extra periodicity, being the interplanar distance d110 of the diamond silicon equal to 0.192 nm. The detailed explanation can be found in the Figures S1− S4 in the Supporting Information. In the literature, the lattice periodicity of 0.33 nm has also been observed in very thin diamond Si nanowires (see, for example, refs 18−21) where it has been attributed to the forbidden 1/3 ⟨422⟩ reflections following.25,26 The explanation to justify the presence of the forbidden fractional reflections of the type 1/3 ⟨422⟩ cannot be applied to our case for two main reasons: (1) According to ref 25 the possibility to see such reflections is limited to “exceptionally thin and flat” samples (t < 30 nm as in the case of refs 18−21). As already mentioned and as shown in Supporting Information Figure S5, we observed those reflections in wires with diameters up to 260 nm and not so flat. (2) The papers that deal with these forbidden reflections (see ref 26 and references therein) insist that they are weak; on the contrary, we always find that they are at least as intense as the common [220] reflections, but in some cases they are more intense, as it can be seen in Supporting Information Figure S5. From the TEM analyses we found that the doping species have a significantly different effect on the crystalline properties of the NWs, thereby causing the variation of the relative abundance of two polymorphs, as shown in the histogram in Figure 3a. Forty wires for each sample have been fully characterized. The onset of the small percentage of the

diamond phase in B-doped NWs compared to the undoped ones is probably related to the stress release caused by the formation of planar defects, that is, twin planes (Figure S2 in the Supporting Information). In Figure 3b, an HRTEM image of a typical wurtzite NW in [0001] zone axis is shown with its fast Fourier transform (FFT) as the inset. Such NWs are surrounded by a 1−3 nm thick amorphous Si oxide layer. In Figure 3c, a P-doped diamond NW is imaged by HRTEM in the [011] projection. The vast majority of the NWs are perfect crystals. If defected, they commonly present twin planes parallel to their longitudinal axis, in agreement with Liu et al.23 (see Figure S1 and S2 in the Supporting Information). Very few diamond NWs present short stacking faults perpendicular to their (111) directions, most likely due to aging, as shown for example in Figure 3c. It must be stressed that the degradation of the hexagonal phase in our doped wurtzite Si NWs is slower than in ref 11. The presence of the wurtzite phase is also supported by micro-Raman scattering analyses. It is well-known that the Raman spectrum of bulk Si presents a sharp peak at a wavenumber of 520 cm−1, corresponding to the transverse optical phonon mode of diamond Si11 (Figure 3d). The Raman spectrum of B-doped NWs exhibits a different line shape from bulk Si, that is, the Si optical phonon peak is observed at 515 cm−1. Also in this case an asymmetric broadening toward the higher wavenumber side is observed (Figure 3d). The asymmetric broadening is ascribed to the Fano effect due to B doping.10,27 It is worth noting that the Si optical phonon peak 5902

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is too largely downshifted with respect to the cubic silicon optical phonon, even taking into account the impact of doping and phonon confinement effects28 of the NWs tip. In P-doped NWs, the Si optical phonon (TO) peak is set at 519 cm−1 with an asymmetric broadening toward the lower wavenumber side. The downshift and the broadening of TO peak can be ascribed to doping and phonon confinement effects28 and to the possible convolution with other peaks. An additional peak at 495 cm−1 also appears. Previous works29−31 report on two main TO peaks around 495 and 517 cm−1 that the authors assign to the hexagonal phase of Si. On the other hand, ref 11 shows a peak at 503 cm−1 that is also ascribed to hexagonal Si. Our TO peaks around 495 and 519 cm−1 are more in agreement with refs 30 and 31 rather than with ref 11. We ascribe the small shift from 517 to 519 cm−1 to the presence of 70% of cubic phase in the P-doped NWs that gives rise to a TO peak at 520 cm−1. Therefore, we conclude that the peak at 515 cm−1 observed from B-doped NWs is due to the wurtzite silicon phase.29−31 The asymmetric broadening of the Si optical phonon peak related to the Fano effect masks the 495 cm−1 peak (to support our findings the Raman spectrum of the undoped NWs is also reported in Figure S7 in the Supporting Information). Additionally, B-doped NWs present a further peak at about 618 cm−1, due to the B local vibrational mode.10 The detection of the B local vibrational mode and Fano broadening is the direct evidence that B atoms are located in substitutional sites of the silicon lattice and are electrically active. This analysis further supports the presence of different dominating phases in the samples doped with different atomic species, consistently with the observations made by HRTEM statistics. In order to determine the optical bandgap of Si NWs, their optical properties are studied by means of absorption (SPS) in Figure 4a and emission spectroscopies (CL) in Figure 4b. The surface photovoltage method is based on the measurement of the surface photovoltage as a function of the photon energy. By changing the photon energy also the optical absorption coefficient, α, and in turn the penetration depth, α−1, varies. For photon energies ranging from 0.8 to 1.4 eV, the penetration depth ranges from 106 to 10 μm in bulk Si.32 Thus, in this spectral range the carrier generation and collection occur mainly within the Si substrate, while for photon energies above 1.4 eV, they occur mainly within the NW mats. It is clear from the comparison shown in Figure 4a that, while P-doped NWs and bulk Si show very similar SPV spectra with an onset at the fundamental band gap of Si (∼1 eV), B-doped and undoped NWs show an additional onset at about 1.6 eV. At these photon energies, the penetration depth is within the NW mats, hence this feature can be safely related to the NWs. Urbach states induced by crystal disorder in the structure give rise to these quite smooth optical transitions, which do not allow for a more precise determination of the optical band gap. For this reason, the error in Figure 4a is around 0.1 eV. Therefore the B-doped and undoped NWs result to have a band gap at 1.6 eV while the P-doped ones have the band gap at 1.0 eV. It is worth noting that undoped and B-doped wurtzite NWs possess a transition at an energy comparable with theoretical prediction of the direct transition at the Γ point.33 Further the band edge is in agreement with experimental reports about wurtzite Si structures, that is, wurtzite bulk Si34 and wurtzite Si microcrystals.29 The energy gap of P-doped NWs is almost the same as that of diamond bulk Si.

Figure 4. (a) Surface photovoltage spectra of B-doped (red triangle), undoped (green dots), P-doped (blue dots) Si NWs, and bulk Si (black square). The slope changes correspond to the estimated optical band gaps for B-doped and undoped Si NWs is 1.6 ± 0.1 eV for Pdoped NWs and bulk Si is 1.1 ± 0.1 eV. For photon energies ranging from 0.8 to 1.4 eV, the penetration depth ranges from 106 to 10 μm in bulk Si.32 Thus, in this spectral range the carrier generation and collection occur mainly within the Si substrate, while for photon energies above 1.5 eV they occur mainly within the NW mats. (b) Comparison among the CL spectra of undoped, B-, and P-doped Si NWs (closed symbols), and bulk Si (open symbols). It must be stressed that the CL emission from bulk silicon is magnified 100 times with respect to the other transitions. This results emphasizes the much higher emission efficiency of WZ NWs.

The luminescence emission properties of undoped, B-, and P-doped NWs are investigated by CL spectroscopic analyses. The advantage of using CL spectroscopy in the characterization of NWs is mainly due to a straightforward information on the band gap nature (see, for example, ref 35) and to the high injection power indispensable in case of materials with low luminescence efficiency like Si (see, for example, ref 36) Gaussian deconvolution procedures of the CL emission spectra of all types of NWs reveal the presence of two bands underneath the main broad emission (Figure 4 b). This value is consistent with the absorption edge found by SPV analyses of B-doped NWs (Figure 4a) and with theoretical predictions of 5903

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the direct transition at the Γ point of wurzite silicon.33 By comparing the CL integrated intensities of the 1.53 eV peak for all the NWs, the B-doped NWs intensity results stronger than those of the undoped and P-doped ones. The undoped NWs have almost the great majority of wurtzite phase but have the same CL integrated intensity of the P-doped ones; further, the NW densities we measured are 8 × 108 and 1012 cm−2 for Bdoped and P-doped mats, respectively. From these observations, it follows that the more intense CL emission of the Bdoped samples can only be due to their larger diameter (see Figure S11 and S12 in the Supporting Information) with respect to the undoped and P-doped NWs. Actually, as shown from Monte Carlo simulations (Figure S10 in the Supporting Information), the shape and dimension of the generationrecombination volume and the CL efficiency depend on the NW diameter. As a consequence, larger NWs give higher CL integrated intensity . The diameters of B-doped NWs exhibit a broad, Gaussianlike, distribution, ranging from 100 to 530 nm. Their mean value is 380 nm with a standard error of ±50 nm. Differently, the diameters of the P-doped NWs exhibit a strongly asymmetric distribution with a sharp peak followed by a long tail, where the diameters range from about 15−160 nm. Their mean diameter is 26 nm with a standard error of ±9 nm. Considering the diameter statistics of all types of NWs, no quantum confinement effect can be expected37 (see Figure S12 in the Supporting Information). The CL band around 1.68 eV is ascribed to SiOx-related luminescence centers, as reported by ref 38 or to amorphous silicon.39 This hypothesis is also supported by our HRTEM observations that show the presence of a few nanometers-thick SiOx outer layer on all types of NW core (Figure 1a, and Figure S1 in the Supporting Information). Assuming a Stokes shift of 60 meV,40 we ascribe the CL emission band at 1.06 eV to the NBE transition of the diamond-like structure of bulk Si (Figure 4b). It is worth noting that in order to obtain the CL emission intensity by the bulk Si comparable with the NWs one, an injection power 2 orders of magnitude higher must be used. Therefore, after renormalizing all the integrated intensities to the injection power, the results are (in arbitrary units) 2750, 2786, 3195, and 15 for undoped, P-, and B-doped NWs, and bulk Si, respectively. This finding demonstrates that wurtzite NWs possess a radiative efficiency 2 orders of magnitude higher than cubic bulk Si. Quantum efficiency analyses are carried out (see Figure S8 in the Supporting Information) to clarify the effect of the wurtzite structure on the optoelectronic behaviors of electrically active Si NWs. Reflectivity studies are also performed on both types of doped NWs and on the underlayer Si substrate after removing the NWs. The B- and P-doped NW mats exhibit a stronger suppression of the optical reflection in respect to the Si underlayer over a wide energy range extending from 0.7 to 2.3 eV (Figure 5). This remarkably low reflectivity is attributable to the multiantireflection behavior of the NWs.41 A further cause of the reflectivity lowering could be the presence of the SiOx outer layer present in all the nanowires (Supporting Information Figure S1). As a matter of fact, it is well-known that boron/diborane assists in decomposing SiH4 to Si. Thus, the difference between the B-doped Si-NW reflectivity and that of P-doped Si-NW one is likely due to the semiamorphous Si overcoat that grows more rapidly on the B-doped Si NWs than on the P-doped. This hypothesis is also supported by the CL results (Figure 3b) on the emission band at 1.68 eV, ascribed to

Figure 5. Reflectivity spectra of B-doped (red line) and P-doped (blue line) Si NWs. The spectrum of the Si underlayer (black line) after removing the NWs from the mats is also reported.

SiOx-related luminescence centers, or amorphous silicon. It is worth stressing that the difference in reflectivity of B- and Pdoped NWs ranges from 2 to 15% in the whole energy range. We suggest that this behavior can be related to the peculiar morphology, density, and average diameter of B-doped NWs (see Figure S11 in the Supporting Information). In fact, it is known that the tapered end of nanowires reduces the reflectivity with respect to nontapered wires. It must be stressed that the reflectivity curve of the silicon underlayer after removing the NWs is a bit lower than that of a perfect commercial silicon substrate due to the residual damage caused by the NW removal. This result clearly demonstrates that the difference in quantum efficiency between B- and Pdoped NWs is affected not only by their reflectivity. In this respect, also the different densities of B-doped and P-doped NWs mats should be taken into account (see Figure S11 in the Supporting Information). Therefore, B-doped NWs have considerable potential for solar cell applications because they possess an extremely lower reflectivity with respect to bulk Si over a wide spectral range in the visible/near-infrared region due to the presence of the wurtzite polymorph. In the following, we propose a simplified model to explain the different effectiveness of boron and phosphorus in retaining the hexagonal phase in our Au-catalyzed Si NWs. Extensive studies have been carried out on the phase changes induced by doping42 on III−V NWs. Differently, to the best of our knowledge there are neither theoretical nor experimental reports on the same topic on wurtzite Si NW. Fontcuberta i Morral et al.11 have recently demonstrated the presence of different polymorphs in undoped Si NWs mats as a consequence of the large surface tension of the Au nanoparticles catalyst during the NW nucleation. 5904

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Notes

It is known that boron and phosphorus have different solubilities in gold with boron keeping the lowest (see refs 43 and 44 where it is reported that the solubility of boron in gold is lower than 5%). Therefore, most likely during the growth of our gold-catalyzed NWs, the surface tension of the Au nanoparticle is not affected by the presence of boron thus resulting in a growth mechanism similar to that of undoped wurzite Si NWs that predominantly are present in the wurtzite phase (Figure 2a). On the contrary, phosphorus has a solubility in gold of 13%45 and therefore the surface tension of the gold catalyst NP can be strongly modified by phosphorus thus inducing a different growth mechanism giving rise to the nucleation of a significant amount of diamond nanowires. Since this is the first experimental report on the influence of doping in retaining the hexagonal phase of wurtzite silicon NWs and because of the lack of any models in the literature on this topic, alternative explanations can be adopted for explaining our findings. For instance, the possibility that at high growth rates the P or B may change the growth kinetics leading to one phase or the other being more favorably stabilized could be taken into account. However, to the best of our knowledge it is not possible to discriminate among the different explanations without experimentally defining the ternary phase diagrams involving silicon, gold, and dopant atoms that are not presently available in the literature. That is however beyond the purpose of the present work. In this study, we prepare the way for p- and n-type doping silicon nanowires while retaining their wurtzite phase. HRTEM studies demonstrate that boron- and phosphorus-doped NWs, respectively, keep about 67 and 28% of wurtzite phase in the mats with the two polymorphs never coexisting in the same NW. Our optical investigations also prove that wurtzite nanowires present a direct transition at the Γ point around 1.5 eV. On the contrary, the emission of diamond nanowires occurs approximately at 1.1 eV. Moreover, a remarkably lower reflectivity in the range from 0.7 to 2.2 eV is reported for Bdoped Si NWs mats, reasonably due to both structural and morphological features. Therefore, we strongly believe that B-doped wurtzite NWs have considerable potential for solar cell applications due to the direct transition and to a low reflectivity.

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The authors declare no competing financial interest.

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ACKNOWLEDGMENTS N.F. was supported by a Funding Program for Next Generation World-Leading Researchers (NEXT Program) in Japan. The research has been partly supported by the Italian-Japanese Progetto di Grande Rilevanza “Nanoscale assEssMent of chEmical and phySical propertIes of advanced nanoStructures (NEMESIS)“ funded by the Italian Ministry of Foreign Affairs. Thanks are due to Prof. G. Calestani for X-ray investigations.

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ASSOCIATED CONTENT

S Supporting Information *

TEM image of a defected cubic wire on the B-doped sample, atomic model of the two structure that could be confused during the TEM observations, HRTEM simulations of the possible structures at different defocus, XRD analysis of Bdoped NWs, Raman analyses of the undoped Si-NWs mat, quantum efficiency of P-doped NWs mats and silicon underlayer, Gaussian deconvolutions of the CL spectra, Monte Carlo simulation for the generation-recombination volume, morphological analyses of the samples, BF-TEM image showing the tapering of the B-doped NWs, three dimensional map of the energy released along the diameter of Si-NWs, schematic for the PCS characterization. This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. 5905

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