Strong Room-Temperature Visible Photoluminescence of Amorphous

May 21, 2018 - considered as a relatively low cost alternative because of its simplicity and .... amorphous domain to amorphous hydrogenated bulk sili...
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Strong room-temperature visible photoluminescence of amorphous Si nanowires prepared by electrodeposition in ionic liquids Shibin Thomas, Jeremy Mallet, Florie Martineau, Hervé Rinnert, and Michael Molinari ACS Photonics, Just Accepted Manuscript • DOI: 10.1021/acsphotonics.8b00208 • Publication Date (Web): 21 May 2018 Downloaded from http://pubs.acs.org on May 21, 2018

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is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

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Figure 1: (a) Schematic representation of the electrochemical setup used for electrodeposition of Si NWs (b) cyclic voltammogram of 0.1 M SiCl4 in Py1,4[TFSI] recorded on Au substrate at 50°C with a scan rate of 10 mV.s-1.

community is also working on the improvements to be done to make the process industrially compatible so as to scale up the deposition for future large scale applications. This templateassisted electrodeposition technique is catalyst free and allows performing the deposition of amorphous Si NWs in the ambient conditions. The NW characteristics are easy to control as their diameters are directly linked to the diameters of the template pores and their length to the thickness of the template. By using this technique, it is possible to grow NWs with diameters from around ten nanometers to few hundreds of nanometers31-33 passivated with a native surface oxide. In this study, the optical properties of the electrodeposited Si NWs with diameters ranging from 110 to 400nm are investigated by PL and fluorescence measurements. In correlation with the NW structural properties, it is shown that the NWs exhibit strong emission in the red with quantum efficiency superior to 25% for the 110nm NWs, depicting the interest of our simple and relatively low-cost growth process to obtain NWs useable for optical applications. This radiative emission does not seem related to radiative defects, but its temperaturedependent and diameter-dependent behaviours are characteristic from band-tail recombination and could be explained by spatial confinement of the carriers within the amorphous NWs passivated by a native surface oxide, the amorphous state being extremely important to get emission at room temperature (RT). While performing the template-assisted electrodeposition of NWs, an important primary step is to confirm the wettability of the nanopores with the electrolytic medium. Unlike in the case of electrodeposition in the aqueous medium, where the viscosity of the electrolyte is negligible, the ionic liquid medium requires special attention because of the bulky nature of their ions. The bulky ions in the IL and the consequent enhancement in the viscosity may have an impact on the dewetting properties of the liquid, especially when used in a nanoporous template. The simplest electrochemical methods to confirm the wettability is to monitor the open circuit potential (OCP) and to perform a cyclic voltammetry (CV) of the electrolyte using the nanoporous template. In addition to giving important information about the dewetting properties, the CV is also used to understand the electrochemical processes occurring at the electrode/electrolyte interface. This knowledge of electrochemical processes is quite important for choosing the appropriate potential for conducting the electrodeposition in the further step. A schematic representation of

the electrochemical setup used for the electrochemical characterizations and the deposition of the Si NWs is shown in figure 1a. Figure 1b shows the CV of 0.1 M SiCl4 in Py1,4[TFSI] recorded on Au substrate at 50°C with a scan rate of 10 mV.s-1. The CV gives three prominent reduction waves in the cathodic scan, marked as C1, C2 and C3 in the figure 1b. The presence of various reduction peaks with significant current intensities is a clear indication of the wetting of nanopores by the IL electrolyte. The broad peak C1, observed around -1.3V is attributed to the reconstruction of Au surface as already reported in similar electrochemical system37. The peaks C2 and C3, observed at -2.1V and -3.0V respectively, are assigned for the bulk reduction of Si onto the working electrode through the nanopores. The origin of the two electrochemical processes corresponding to the peaks C2 and C3 might be the reduction of two different complexes of Si4+ ions which necessitates two different over-potentials to get reduced. By polarizing the working electrode in the whole potential range from -2.0V to 3.0V, Si electrodeposition can be achieved. The plateau observed from -2.9V till -2.1V on the anodic scan also confirms the stable diffusion limited process occurring at this potential range leading to the reduction of Si. The exponential increase in reduction current observed after -3.2V is due to the breakdown of the IL and the irreversible reduction of (Py1,4)+ cation. The CV was used to identify the potential at which the Si reduction occurs, which could be used to obtain Si NWs by electrodeposition. The Si NWs were then potentiostatically electrodeposited onto the PC membranes with 110nm, 200nm and 400nm pore diameter and with a thickness of 20µm by applying -2.8V at 50°C. Figure 2a, 2b and 2c shows representative typical SEM images of 400, 200 and 110nm Si NWs respectively electrodeposited from the IL. The images display dense networks of Si NWs with quite uniform diameters and length for each condition. For PC templates with different diameters, the deposited NWs are characterized by a fixed diameter and length up to few micrometers, which perfectly corresponds to the diameter and length of the PC membranes used. Hence the technique is quite useful for obtaining large quantities of NWs with uniform length and diameters. As can be seen from the SEM images, the obtained NWs keeps their uniformity in terms of shape and size even after undergoing the aggressive cleaning protocols utilizing chloroform, indicating their robust nature. Raman spectroscopy was employed to get information

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Figure 2: Representative typical top view SEM images of 400 (a), 200 (b) and 110nm (c) Si NWs electrodeposited from the IL. (d) The Raman spectrum measured from 110nm Si NWs. (e) EDS spectrum recorded from 110nm Si NWs.

about the presence of Si and the amorphous or crystalline nature of the NWs. Raman spectroscopy is a widely accepted technique used to distinguish between the crystalline (c-Si) and amorphous (a-Si) forms of Si. It is known from the literature that in the Raman spectrum, the c-Si shows a very thin band at 520 cm-1, which is due to the scattering of first-order transverse optical phonon38. When it comes to the a-Si39, the selection rules are broken and it exhibits two broad bands at 150 cm-1 and 480 cm-1 associated with the Transverse Acoustic (TA) and Transverse Optic (TO) phonons respectively. A typical Raman spectrum obtained from as-deposited 110nm Si NWs is shown in figure 2d. It is evident that the peak at 520 cm-1 is absent in the current Raman spectrum and hence we rule out the possibility of c-Si in the NWs. Instead, two broad peaks can be identified from the spectrum, one at 150 cm-1 and the other at 480 cm-1, which well corresponds to the Raman shifts observed for a-Si. Therefore, it is inferred from the Raman spectrum that the as-deposited Si NWs are amorphous in nature, whatever their diameters be, we always get similar spectrum. In order to get information on the chemical composition of the NWs, EDS analysis was performed. A typical EDS spectrum obtained from a bunch of 110nm Si NWs deposited on a Ni TEM grid (figure 2e) exhibits signals corresponding to Si, O, C and Ni and similar spectra are obtained from all the NWs with different diameters. The Ni and C signals are arising from the Ni grid used as a substrate to collect the NWs. The strong signature of Si in the EDS spectrum confirms that the deposited NWs are composed of pure Si mainly. The presence of oxygen could be related to the unavoidable oxidation of the silicon surface during the washing process or when exposed to the ambient atmosphere. No other impurities were detected from the NWs, indicating the compositional quality of the deposited NWs. EDS mapping shows a uniform content of silicon and oxygen over the NWs as seen

in figure 3 (top) along the length of the NWs (low magnification images) with low composition disparities (higher magnification images in the inset of the figures). The oxygen content in the NWs disappears after the HF etching, as evident from the EDS mapping (figure 3 bottom). EDS analysis performed on isolated single NWs did not evidence the presence of impurities within the NWs in the limit of the sensibility of the techniques, even if long counting times were used to try to improve the detection of possible low concentration elements. This absence of impurities, possibly coming from the ionic liquid, has also been confirmed by X-Ray Photoelectron Spectroscopy analysis (XPS) (not shown here) performed on a bunch of nanowires, and thus giving only average results compared to the mapping obtained by EDS. These purity of the silicon structures (thin films, NWs, nanograins) has also been confirmed by the work of our group40 or other groups41-42 using Secondary Ion Mass Spectroscopy, XPS or Scanning Tunneling Microscopy. To study the optical emission properties of the NWs, PL and fluorescence experiments were then carried out. The room temperature PL emission spectra obtained from the NWs with different diameters are shown in figure 4a. The three samples exhibit the same broad band located around 480nm and a second band at higher wavelength in the red (>780nm) which depends on the NW diameters. For the 110nm NWs, the intensity of the 480nm band is weak compared to the peak at 780nm, and for the 200nm NWs, the 480nm band intensity is still inferior to the one at higher wavelength which is shifted to 830nm. The 400nm NWs have a different behaviour as the band at higher wavelength is shifted to 860nm but with a very weak intensity compared to the band at 480nm. The broad peak at 480 nm has already been observed in other silicon-based systems and it is often attributed to the defects

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Figure 3:(top) Low magnification TEM image (left) and the EDS elemental map of Si and O (right) from as-deposited Si NWs. The inset shows the high magnification TEM image and the corresponding EDS elemental map of Si and O. (bottom) Low magnification TEM image of Si NWs after HF treatment (left) and the corresponding Si and O elemental mapping (right).

in silicon oxide (SiOx)27, 43-44. As the silicon and oxygen EDS mapping (figure 3 top) shows that the grown NWs undergoes oxidation, this band can be coherent with oxygen defects. As shown in figure 4b, the temperature-dependence for the emission wavelength of the 110nm NWs is consistent with a defect-like emission as the wavelength remains constant with the temperature (no PL shift) and the line shape is not changing. As for the intensity, the strong decrease between RT and 4K is also not relevant for the behavior of a semiconductor intrinsic emission. For the higher emission red band, such band has already been observed between 700 and 900nm in different amorphous silicon systems1, 22, 45 from p-Si containing some amorphous domain to amorphous hydrogenated bulk silicon46 and related films47 or amorphous silicon nanocrystals48. Many models have been used to explain this emission in the red from quantum confinement49-51 to oxide-related extrinsic defects52, surface states such as hydrides or polysilane53-54 or band-tail recombination55-57. In our case, the NWs diameters are well above the excitonic Bohr radius of silicon and then a quantum confinement mechanism could not be the proper model to explain the emission. Regarding the temperature-dependence of this band for the 110nm NWs (figure 4b), it is clearly not the behaviour of a radiative defect as the band is redshifted with the temperature increase and at the same time, the intensity decrease is very limited. Such redshift with temperature has already been observed and is more explainable by a bandtail radiative recombination between the deepest energy accessible conduction and valence-band states via tunneling. But at the same time, even if we are not in a quantum confinement regime, a size effect is also observed with a redshift from 780 to 880nm when the NW diameters are increasing from 110 to 400nm. To quantify the difference in PL intensity with the NWs diameters and as in the different sets of NWs, especially the 110nm NWs were brightly emitting at RT, we then performed fluorescence measurements to have quantum yield (QY) values. It appears that the 110nm, 200nm, and 400nm NWs have a QY of 28±8%, 8±5%, and 0.5±0.3%. The QY for the 110 and 200nm are particularly high for Si NWs either

amorphous or crystalline. This variation of the PL intensity and the PL wavelength with temperature and diameters, and the strong QY is explained by a band-tail recombination mechanism coupled to a spatial confinement of the carriers within the NWs, a model used to explain the PL observed in aSi:H nanostructures which can exhibit high QY. The observed PL from the a-Si NWs could be originating due to the spatial confinement effects as previously observed and modelled in similar amorphous silicon systems56-58. In amorphous silicon systems, the electron and hole wavefunctions are highly localized due to structural disorders and can radiatively recombine. In bulk a-Si films, the efficiency of the emission will be quite weak because of the possibility for the electrons and holes to recombine via non-radiative processes through defects. In the quantum confinement model, which is commonly used to explain the luminescence from NWs with very low (sub-10nm) diameters59-60, the light emission is resulting from the modification of the band gaps due to the confinement of the carriers. When the NW diameter decreases and reaches an optimal limit known as the exciton Bohr radius, which is 4.9nm for Si, the energy bands of the NWs turns to be discrete and its band gap becomes dependent on the NW diameters61. As the dimensions of our NWs are well above the exciton Bohr radius, and regarding the evolution of the PL intensity with temperature, the quantum confinement model could not be applied to explain the observed PL emission properties. This model fails to resolve the large variations of PL energy and intensity observed for the electrodeposited Si NWs as far as the lifetime in the range of the µs (fig. 4(b) and (c)). On the other hand, the spatial confinement can be understood in a statistical viewpoint where the probability of the electron-hole pair to encounter a defect state before undergoing radiative recombination is lowered in a confined volume. In such confined systems, the defect mediated non-radiative recombination processes, which are dominant in bulk a-Si at RT, are essentially reduced. In the spatial confinement model, the photo-generated electron hole pairs thermalizes rapidly to the lowest energy states

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Figure 4: (a) The PL emission spectra at RT from the electrodeposited Si NWs with 110, 200 and 400nm diameter. The table at the inset shows the quantum yield estimated from each NW. (b) The temperature dependence of the PL emission bands obtained from the NWs. The solid lines represent the PL intensities while the dashed lines the PL wavelength. (c) PL decay curve for 110nm Si NWs. (d) The comparison of normalized PL spectra of the as-deposited NWs (black line), the NWs after HF treatment (blue line) and the NWs after high temperature annealing. The PL experiments have been performed at room temperature for the amorphous NWs and at 4K for the crystallized NWs.

within a capture radius . The capture radius, , is given by the equation below.  

 

ln  1

Where  is the localization radius, is the escape frequency from localization states and  is the radiative life time. The capture volume ( for electron-hole pair is given by, 4    2 3 And the radiative efficiency for an electron-hole pair is given by      3 Where  is the defect density. It is clear that when the size of the domains are lower than the , the electron-hole pair undergoes spatial confinement and consequently the capture volume decreases. This results in an increase in the radiative efficiency according to equation 2. Also, as the capture volume decreases, the probability of encountering a defect state by the electron-hole pairs decreases because of the relatively low number of defect states in a confined volume. On the other hand, this increases the radiative recombination efficiency. This model has been successfully employed in the past to explain the visible PL mechanism from porous a-Si nanostructures1, 58, 62 and PL from c-Si NWs with large diameters30, 63. Our results are well described by this spatial confinement model. Regarding the PL intensity of the 780nm peak which shows only a small variation between 4K and the ambient temperature, such weak temperature dependence of the PL

intensity of a-Si NWs are totally different from the PL properties of bulk a-Si, where a strong thermal quenching is generally observed due to the non-radiative recombination centers related to structural disorders and dangling bonds64. In comparison with the bulk a-Si, a strong improvement of the thermal quenching of the PL intensity is obtained by fabricating the a-Si NWs. The temperature dependence of the PL intensity is controlled by the thermally activated process of electronhole pairs localized in the band-tail states. The spreading of the band tail states in a-Si NWs, achieved via the nanostructuration of the material, decreases the thermal emission rates of carriers in the band tail states and reduces the thermal quenching of the PL intensity65. As for the PL wavelength of the NWs, the model predicts emission peak energies for a 1-D aSi nanostructure which are coherent with our values. The shift of the emission wavelengths with the NW diameters is not linked to the quantum confinement mechanism as the diameters are too large here, but such red shift is qualitatively also predicted by the spatial confinement model while the NW diameters are increasing. As for the intensity, smaller diameters will limit the probability for the carriers to be captured by defects. It can then explain the decreasing of the QY with the increase of the diameters as in the bigger NWs, the number of defects will surely be higher and then the probability for the carriers to radiatively recombine will be quite small. From a quantitative point of view, if compared with the model of Estes et al., the high QY value and the observed emission wavelength for our NWs are more suitable with smaller amorphous domains. Nevertheless, it can be explained by the fact that our fabrication process is not a crystallographic growth process and that small variation in silicon density or defect

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density could appear along the length of the NWs. Since the electrodeposition is a bottom-up nanofabrication method, the deposited a-Si NWs are surely composed of amorphous silicon domains with a dimension of a few tens of nanometers with different densities rather than a continuous structure. Then the capture radius will be shortened within the NWs leading to a higher QY and explaining the reduced temperature dependence for the PL intensity. Our explanation is also supported by time- resolved PL performed on the 110nm NWs. A typical decay curve at RT is shown in figure 4c and could be fitted by a double-exponential decay curve. The PL decay profile consists of two components at 0.780±0.140 µs and 2.9±0.3 µs. Similar lifetimes in the order of the µs have already been observed in amorphous silicon films65-67. The measured values in the order of the µs are compatible with the spatial confinement model56-57 (and not with the quantum confinement mechanism) with decay in the order of 10-6-10-7s if the confinement diameter is in the order of a few tens of nanometers, as we stated for the structure of our NWs. The two different, but very close, lifetimes could be due to the intrinsic disorder in the amorphous NWs because the growth process involves some oxygen-passivated domains and some pure amorphous domains as in the study of Tsushima et al.67 where they attributed a µs lifetime in a-Si:H films and nanostructures to recombination of exciton in confined hydrogen-free nanostructures. Regarding all these different PL results, we can then assume that the red emission and the high QY for the 110nm NWs could be ascribed to a band-tail recombination mechanism with spatial confinement of the carriers limiting the non-radiative processes. In order to be sure of the origin of the two emission bands observed for the NWs, an HF etching process was carried out on the 110nm NWs, aiming to remove the silicon oxide shell from the surface of the NWs and replace the SiOx surface coverage by a Si-H passivation. After 5 minutes of HF treatment, the EDS mapping clearly shows the absence of oxygen signal in a typical 110nm NWs (figure 3 bottom right) confirming that the oxygen detected for the as-deposited NWs is coming from the oxidation of the surface and the creation of a SiOx shell. PL measurements (figure 4d) carried out on the NWs show the disappearance of the 480nm band but the red band is still present, slightly blueshifted and broader compared to the 780nm band of the as-deposited NWs. This confirms that the 480nm band is related to oxide defects which disappeared with the HF treatments. As for the red band, its behaviour is more coherent with a band-to-band recombination mechanism. The slight blueshift could be due to a slight decrease in the diameter of the Si NWs during the etching process, or rather to the fact that the Si-O bonds are replaced by Si-H bonds, which do not possess the same energy and then induces a slight difference in the energy of the band tails leading to a slightly larger band-gap. At the same time, the HF etching could create heterogeneities and irregularities on the NW surface and hence could increase the disorder within the band tails of the a-Si leading to the broadening of the emission band. Then, in order to confirm that the red emission and the high QY is due to the amorphous nature of the Si, the as-deposited 110nm NWs were further annealed at 650°C to transform the amorphous phase to crystalline phase as previously evidenced31. After the crystallization, the band at 480nm could still be observed due to the oxidized surface, however, the band at 780 nm was not detected in the PL but replaced by a narrow band at 1180nm (1.06 eV) close to the band-gap of the

crystalline silicon (figure 4d), which is almost non-detectable at RT but observable at 4K. Such band has already been observed in crystalline Si NWs of similar diameters30 and has been attributed to a band-to-band recombination combined to spatial confinement of the carriers within the NWs. As for its intensity, the crystallization process obtained by annealing should imply the appearance of numerous interface defects limiting the radiative emission at RT. This disappearance of the 780nm band with the Si crystallization combined with the appearance of the band at 1180nm close to the c-Si band-gap confirmed that the amorphous state of the Si NWs is directly responsible of the red emission observed with our electrodeposited NWs. In summary, amorphous silicon nanowires passivated by a native surface oxide and with diameters from 110 to 400nm were obtained by a template-assisted electrodeposition method from the ionic liquids. The electrodeposition produces a dense network of Si NWs with quite uniform diameter and length. The NWs are found to be clean and possess good structural and compositional quality. The Raman spectroscopic measurements reveal that the as-deposited Si NWs are amorphous in nature. The Si NWs show strong PL at room temperature in the visible wavelengths, characterized by the presence of two emission bands identified at 480nm and 780/800nm depending of the NW diameters. While the 480nm luminescence is originating from the defect bands in SiOx, the strong and efficient red band should result from intrinsic radiative recombination and spatial confinement of the carriers in a-Si NWs well passivated by an oxide layer. The temperature dependence of the emission wavelengths and the PL intensity also confirms the origin of the luminescence bands from the amorphous states in Si NWs leading to QY >25% for the 110nm NWs passivated by a native surface oxide. The present study elucidates a new alternative growth technique for obtaining luminescent Si NWs with high QY with potential applications in Si-based optoelectronics. Experimental Section The Si NWs were prepared in an electrolytic bath containing 1-butyl-1-methylpyrrolidinium bis(triflouromethanesulfonyl)imide (Py1,4[TFSI]) (solvionic S.A) and SiCl4 (0.1M) (Aldrich). The preparation of the electrolyte and the electrodeposition was performed inside an argon filled glove box with moisture and O2 content less than 1 ppm. Commercially available polycarbonate (PC) membranes with 110, 200, and 400nm pore diameters served as the template for the NW deposition. Prior to the deposition, a thin layer of gold (approximately 200nm) was sputtered on one side of the PC membrane to make the electrical contact during the electrodeposition. The electrodeposition was performed using a three electrode based electrochemical cell. Platinum wires were used as both the reference electrode and the counter electrode. Utmost care was given for the cleaning and drying of the cell and the electrodes before conducting the electrochemical experiments. All the electrochemical measurements were done using a Voltalab PGZ 100 potentiostat/galvanostat controlled by Voltamaster 4 software. After the electrodeposition, the Au back electrode on the PC membrane was removed fully by ultra-sonication in ethanol for few minutes. The NWs were then separated from the template by dissolving the PC membrane in distilled chloroform. The suspension was then centrifuged and washed multiple times with the chloroform. The clean NWs were collected either onto Si (100) substrates for scanning electron microscopy (SEM) and PL measurements, to a Ni grid for transmission

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electron microscopy (TEM) experiments and directly put in a proper glass cuvette for fluorescence measurements. For each experiment, multiple samples coming from different batches were analysed to ensure the repeatability of the results. The structural, morphological and compositional characterization of the Si NWs was made by SEM (HITACHI S-3400N) and TEM (Philips CM30) equipped with Energy Dispersive X-ray (EDS) spectrometer. Raman measurements were carried out with a multichannel Jobin Yvon T64000 Raman spectrometer equipped with a 1800 grooves mm-1 grating. For the steady-state PL experiments, the samples were excited by a 30 mW He–Cd laser using the 325 nm line. The PL signal was collected with a liquid nitrogen cooled Si detector and by a monochromator equipped with a 600 grooves mm-1 grating and by an InP/InGaAs photomultiplier tube cooled at 190K, with a detection range in the 600–1700 nm range. PL experiments were performed at temperatures between 10K and 300 K in a cryostat cooled with liquid helium. The quantum yield measurements were performed with a Horiba-Jobin Yvon Fluorolog 3 System with a Xenon lamp as a source. For the quantum yield, classical fluorescence measurement protocols were performed following the procedures setup by the European and American standardization labs using IR 125 as a standard sample68-69 for the relative QY measurements. These measurements were confirmed using an integrative sphere for absolute QY values.

AUTHOR INFORMATION Corresponding Author * E-mail: [email protected]

Notes The authors declare no competing financial interests.

AUTHOR CONTRIBUTIONS S.T., J.M., H.R., and M.M. wrote the manuscript. S.T., J.M., H.R. and M.M. participated to the scientific discussions. S.T., F.M. and J.M. elaborated the nanowires. H.R. and M.M. performed the optical experiments. S.T. and M.M. treated the data for the manuscript. J.M. supervised the electrodeposition part and M.M. was the coordinator of the project. All authors read and approved the final manuscript.

ACKNOWLEDGMENT The authors would like to acknowledge Region ChampagneArdenne, the DRRT Champagne-Ardenne and the European FEDER for the financial support through the Synapse project and their support of the Nano’mat platform. They also acknowledge Laurence Wortham for the technical assistance with TEM meaurements and Nicolas Bercu for the image treatment.

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66. Wehrspohn, R. B.; Chazalviel, J.-N.; Ozanam, F.; Solomon, I., Spatial versus quantum confinement in porous amorphous silicon nanostructures. The European Physical Journal B Condensed Matter and Complex Systems 1999, 8 (2), 179-193. 67. Tsushima, K.; Nakata, H.; Monji, K.; Deki, H.; Murayama, K., Luminescence decay in hydrogenated amorphous silicon and silicon nanostructures. Journal of Non-Crystalline Solids 2012, 358 (17), 2090-2095. 68. Rurack, K.; Spieles, M., Fluorescence quantum yields of a series of red and near-infrared dyes emitting at 600− 1000 nm. Analytical Chemistry 2011, 83 (4), 1232-1242. 69. Würth, C.; Grabolle, M.; Pauli, J.; Spieles, M.; ReschGenger, U., Relative and absolute determination of fluorescence quantum yields of transparent samples. Nature protocols 2013, 8 (8), 1535.

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For Table of Content only “Strong room-temperature visible photoluminescence of amorphous Si nanowires prepared by electrodeposition in ionic liquids” by Shibin Thomas, Jeremy Mallet, Florie Martineau, Hervé Rinnert, and Michael Molinari Amorphous Silicon nanowires prepared with electrodeposition in ionic liquid and with diameters ranging from 110 to 400nm exhibit a strong photoluminescence in the red with quantum yield up to 25%, thanks to a spatial confinement mechanism.

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