SnxGe1âx Alloy Nanocrystals: A First Step toward Solution-Processed

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SnGe Alloy Nanocrystals: A First Step Toward Solution-Processed Group IV Photovoltaics Karthik Ramasamy, Paul G Kotula, Andrew F. Fidler, Michael T. Brumbach, Jeffrey M. Pietryga, and Sergei A. Ivanov Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.5b01041 • Publication Date (Web): 16 Jun 2015 Downloaded from http://pubs.acs.org on June 19, 2015

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Chemistry of Materials

SnxGe1-x Alloy Nanocrystals: A First Step Toward Solution-Processed Group IV Photovoltaics Karthik Ramasamy,a Paul G. Kotula,bAndrew F. Fidler,c Michael T. Brumbach,b Jeffrey M. Pietryga,c Sergei A. Ivanova* a

Center for Integrated Nanotechnology, Los Alamos National Laboratory, Los Alamos, NM, 87545, USA b Materials Characterization Department, Sandia National Laboratory, Albuquerque, NM, 87185, USA c Center for Advanced Solar Photophysics, Los Alamos National Laboratory, Los Alamos, NM, 87545 USA Email: [email protected]

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Abstract: Non-toxic, sustainable and cost-effective, Group IV materials are attractive for a broad range of electronic and opto-electronic applications, although the indirect principal band gaps of silicon and germanium (Ge) present complications that impact device design and cost. Previous studies have shown that the band structures of these materials can be modified by the influence of quantum confinement in nanostructures, or by alloying with tin (Sn) in metastable thin films; to date, neither method has produced a material with a direct band gap of appropriate energy for application in, e.g., efficient solar photovoltaics. We have developed a facile colloidal method for the synthesis of size-controlled, homogeneous SnxGe1-x alloy nanocrystals (NCs) with remarkably high tin concentration (x up to 0.42). We demonstrate that NCs of the same size exhibit a pronounced, systematic red-shift in the optical band gap, and a significant increase in molar absorptivity, with increasing Sn-content, and a measurable photoluminescence was observed from NCs with high contents. The indications of at least partial direct-gap character in these NCs, combined with their broad tunability throughout the infrared, suggest their promise for use in solution-processed solar cells.


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Introduction: The Group IV semiconductors silicon (Si) and germanium (Ge) are of tremendous technological importance, as they are the most common building blocks of modern electronic devices. Their relatively small band gaps (1.1 eV and 0.7 eV, respectively) have given them a particularly large presence in infrared (IR) detectors1 and solar cells,2-3 despite the complicating factor that the indirect nature of their principal band gaps strongly reduces their ability to absorb light. This property comes with important consequences to the manufacture of these devices. For instance, the active absorbing layer of Si photovoltaic (PV) cells must be 100s of µm thick,4 1-2 orders of magnitude thicker than in cells based on direct-gap materials like CdTe 4 or GaAs.5 Efficient collection of the photogenerated charge carriers over such distances places stringent purity demands on the Si in PVs that account for the majority of the high manufacturing cost of the entire device. The use of semiconductor nanocrystals (NCs) in PVs has been largely motivated by the need to reduce their cost. Toward this, NCs offer at least two important advantages: 1) they are amenable toward low-cost, solution-based processing; and 2) they can exhibit unique physical phenomena, due to quantum confinement, that can enhance the efficiency of NC-based devices.6,7 Examples of this latter category include size-tunable band gaps that allow narrower-gap materials to be tuned into the Shockley-Queisser ideal single-junction absorber region of ~0.9-1.3eV,8 and the phenomenon of “carrier multiplication,9” through which single ultraviolet photons can create more than one electron-hole pair, boosting the current collected by a device.10 To date, NC PV efforts have explored a range of materials, including II-VI,11-12 IV-VI,13-14 and I-III-VI 15-16 compound semiconductors, but have largely ignored group IV materials, despite attractive properties such as low toxicity and high abundance. Given that this is mostly due, again, to the overall lower light-absorbing efficiency of these materials, methods for enhancing the absorption of these materials could stimulate widespread investigations of Group IV NCs for PVs and other optoelectronic applications. Numerous theoretical studies suggest that the band structures of Si and Ge can be modified to reduce the energy difference between the first direct and indirect transitions, potentially to the point of “cross-over,” either through the influence of quantum confinement17 or by alloying with tin (Sn).18-20 Indeed, there have been numerous reports of band-edge photoluminescence (PL), an indicator of at least partial direct-gap behavior, in Si21 and Ge22 NCs in the strong quantum confinement regime, although these studies imply that a confinement-induced transition to fully



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direct behavior would be achieved only in materials with band gaps too wide for solar harvesting. Experimental evidence of such behavior in Sn-alloys of either Si or Ge has been more limited, as achieving the requisite Sn concentrations in bulk alloys is very difficult due to the alloying-induced lattice strain and much higher cohesion energies of Si and Ge compared to that of Sn.23 Nonetheless, using molecular beam epitaxy – a relatively costly approach for productionscale fabrication – others have demonstrated that thin films of Ge- and Si-based alloys with Sn appear to show progressively more direct behavior as the amount of Sn increases over the experimentally accessible range of 0-15%.24,25 In contrast to Ge and particularly Si NCs, the resultant alloy films generally possessed too narrow of a band gap (0.27-0.53 eV).25,26 for efficient PV applications. Recently, SnxGe1-x NCs of even higher Sn content, embedded within Ge matrix, has been formed via annealing of metastable epitaxial layer of GeSn solid solution deposited between two Ge layers.27 However, optical and electronic properties of these NCs were not reported. Finally and most intriguingly, low-temperature IR photoluminescence (and even lasing) from an epitaxial layer of GeSn alloy with 13% Sn grown on Si substrate has been attributed to directgap behavior of the material.26 An attractive possibility, clearly, is to harness these two effects in a single material, especially one that retains the advantages of solution processibility (e.g., amenability toward screen printing, spray deposition and other fabrication techniques suitable for roll-to-roll processing or the coating of surfaces of complex morphology). Towards that goal, here, we report the colloidal synthesis of a series of composition-controlled GexSn1-x alloy nanoparticles with Sn concentrations up to 42%. The preparation of such alloys in the nanocrystalline form addresses both drawbacks of the bulk alloys – allowing effective relaxation of the lattice strain created by the introduction of Sn into the Ge lattices, and (via the quantum confinement effect) potentially reducing the amount of Sn needed to achieve direct behavior and widening the final material’s band gap into the solarrelevant region. We show that the introduction of Sn into Ge results in a material system with optical band-gap tunability through a wide spectral region in the IR and, further, that NCs with higher Sn concentrations can exhibit measurable PL, a potential sign of partial direct-gap behavior. As this manuscript was under review, the synthesis of SnGe alloy NCs with tin content up to 27% and with absorption onsets also in the IR, was reported.28 Our analysis of the reaction mechanism suggests that among other differences, our use of lower reaction temperatures allows us to achieve better control over size and size-distribution, more regular particle shape and high-


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er Sn concentrations. Moreover, a more rigorous exclusion of air during physical and spectroscopic characterization of our NCs largely precluded the previously observed surface oxidation, removing a potential source of uncertainty in our conclusions regarding crystal structure and optical properties. █ Experimental

Section

Materials All reagents were used as received, and the solvents were dried over molecular sieves and purged with high purity argon for 30 minutes before use. Germanium (II) diiodide (GeI2), was purchased from Gelest Inc. Tin (II) bis(trimethylsilyl)amide (Sn(HMDS)2, ≥ 99.0%), and oleylamine (OLA, ≥80-90.0%) were obtained from Sigma Aldrich Chemical Co and Acros Organics. 2M n-butyllithium solution in cyclohexane, anhydrous hexane and anhydrous methanol were purchased from Sigma Aldrich Chemical Co. Characterization Routine transmission electron microscopy (TEM) imaging was performed using a FEITecnai, 300 kV transmission electron microscope equipped with a CCD camera for scanning transmission electron microscopy (STEM), a high-angle annular dark field (HAADF) detector, and energy-dispersive Scanning transmission electron microscopy X-ray (EDX) spectroscopy. Scanning transmission electron microscopy (STEM) high-angle annular dark field imaging and EDX elemental mapping were performed using and FEI Company Titan G2 80-200 TEM/STEM operated at 200kV and equipped with the X-FEG (an ultra-stable high-brightness Schottky FEG source), Super-X EDX detector system (4 windowless silicon drift detectors with a combined solid angle of 0.7 sr) and a spherical aberration corrector (CEOS DCOR) on the probe-forming optics. TEM image non-linear processing - Fourier filtering - was carried out using Gatan Digital Micrograph version 3.4. Scanning electron microscope (SEM) imaging was carried out using a Focused Ion Beam Scanning Electron Microscope (FIB-SEM, Carl Zeiss) equipped with EDX. Powder X-ray diffraction (XRD) patterns and small angle X-ray scattering (SAXS) were recorded on a Rigaku Smart Lab II instrument equipped with Cu Kα radiation source operated at 40 kV and 44 mA. Samples for powder XRD measurements were covered with Paratone-N crystallographic oil to prevent excessive oxidation and deposited on the horizontal glass slide. Samples for SAXS were dissolved in hexane and the solution was loaded into 0.7mm boron glass ca-



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pillaries for measurements. PDXL v2.0 and NANO-Solver v3.5 software packages from Rigaku were used to process the powder XRD and SAXS data, respectively. SAXS data were modeled as Gamma distribution of spherical particles with normalized dispersion, , referred to as ‘size distribution’ throughout the rest of the text: [%]=(⟨ R2⟩/⟨R⟩2)1/2100%. XPS measurements was performed using a Kratos Axis Ultra DLD instrument with base pressures less than 5 · 10-9 Torr. Samples for XPS were mounted onto a sample holder in a glove box and briefly (< 5 sec.) exposed to air in transfer from the glove box to the instrument. XPS was performed with a monochromatic Al Kα (1486.7 eV) source operated at 150 W with an elliptical spot size of 300 x 700 microns. Survey spectra were recorded with an 80 eV pass energy, 600-700 meV step sizes, and 100 ms dwell times. High resolution spectra were recorded with a 20 eV pass energy, 50 meV step sizes, and 100 ms dwell times. Charge neutralization was used for all samples to reduce differential charging effects. Data processing was performed with CasaXPS Version 2.3.15. Spectra were adjusted to align the C 1s peak to 284.6 eV. Sn 3d peaks and Ge 3d peaks have been normalized and smoothed to allow for direct comparison between samples. A linear background was applied to the Sn 3d 5/2 peak and the Ge 3d peak. Absorption measurements were carried out using a Cary 6000i dual-beam UV-Vis-NIR spectrometer. PL spectra were taken using a homebuilt setup: excitation from an 808 nm laser was mechanically chopped, and emission was analyzed using a grating monochromator and a LN2-chilled InSb detector using lock-in amplification. Proton decoupled Anasazi EFT 90MHz instrument with the reference

119 119

Sn NMR spectra were acquired using

Sn frequency (0 ppm) set at 33.66574

MHz. Samples were dissolved in pure oleylamine. The acquisition time of 5000 scans, the relaxation delay, and the pulse width were set to 0.14ms, 0.20ms, and 10 s, respectively. TGA/DSC analysis was conducted using Netzsch STA 449 F1 Jupiter simultaneous TGA/DSC instrument using alumina crucibles and dynamic nitrogen atmosphere with total flow of 200mL per min over the sample. Convection Heating Synthesis of SnxGe1-x Nanocrystals via Self-Reduction In a typical reaction, GeI2 (0.2 mmol) was added to a 3-neck round bottom flask containing 5 mL of oleylamine at room temperature. The contents were degassed under vacuum and back filled with N2 for three times and a varied amount of Sn(HMDS)2 was added. The flask was heated to a target growth temperature (between 200oC and 280oC) at which the reaction mixture was held for a period of up to 2 hours before the reaction was quenched by the removal of the


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reaction flask from the heating mantle. The NCs were collected via precipitation by addition of methanol, followed by centrifugation and redispersal in hexane. Convection Heating Synthesis of SnxGe1-x Nanocrystals by Co-Reduction Under the same conditions described above, 60 L of 2M n-BuLi cyclohexane solution were injected into the reaction mixture as soon as the flask reached the target growth temperature. Growth, quenching and NC collection were carried out as above. Microwave Heating Synthesis of SnxGe1-x Nanocrystals via Self-Reduction At room temperature, GeI2 (0.2 mmol) was dissolved in 5 mL of oleylamine by sonication in a 10 mL glass test tube. The contents of the test tube were three times degassed under vacuum followed by back filling with N2. After that, a varied amount of Sn(HMDS)2 in OLA was added. The tube was heated to a target temperature (200-230 oC, typically reached in 15-20 min at 300W) using a CEM Discover microwave reactor (2.45 GHz) under N2 atmosphere. The tube was held at the growth temperature for a period of up to 2 hours before the reaction was quenched by cooling the tube to 55-60 oC. The NCs are collected via precipitation by addition of methanol, followed by centrifugation and redispersal in hexane. Microwave Heating Synthesis of SnxGe1-x Nanocrystals via Co-Reduction Under the same conditions described above, 60 L of 2M n-BuLi cyclohexane solution were injected into the reaction mixture as soon as the tube contents reached the target growth temperature. Growth, quenching and NC collection were carried out as above. Determination of NCs extinction coefficient The stock solution of NCs was prepared by dispersing purified nanocrystals from 0.2 mmol GeI2 reaction in 1 mL of hexane. 50 μL of this solution were diluted to 1 mL using hexane for absorption measurement. To estimate the concentration of nanocrystal dispersion used for absorption measurement, TGA measurement was carried out using 250 μL of the stock solution. The solution was placed into alumina crucible and slowly evaporated to avoid any loss of the sample due to splashing. Final dry was done in vacuum for 15 minutes. In TGA measurements, the sample was held at room temperature for 20 min before it was heated to 500 oC at the rate 15oC/min, after which the sample again was held for 20 min at the final temperature to ensure the complete removal of organic phase. The final mass of the sample was assigned to the Sn xGe1-x



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alloy, which was confirmed by XRD. The molar extinction co-efficient, ε, was determined using the formula: (1) where mTGA is the final mass of the sample in mg; VTGA is the volume of the stock solution used for TGA measurements in μL; VABS is the volume of the stock solution (in μL) that was diluted to 1mL to obtain absorption A; R is the radius of NCs obtained from SAXS measurements in nm, and d is the bulk density (in g/cm3) of the alloy, assuming a linear dependence on composition determined by XRD before the measurement. █Results

and Discussion

Phase Pure Synthesis of SnxGe1-x Alloy Nanocrystals. The contrasting natures of Ge and Sn raise many challenges to the colloidal synthesis of SnxGe1-x NCs. Ge requires a high crystallization temperature due to its covalent bonding nature, but tin melts above 230 °C. While Ge typically crystallizes in a diamond cubic structure, the cubic form of Sn is stable only below 13 °C. Perhaps it is not surprising, then, that the two elements are essentially immiscible (i.e., Sn is only soluble to ~1% in a Ge lattice). Nonetheless, the synthetic methodology developed in this effort is inspired by that used for pure Ge NCs.27 Considering the above complications, we have investigated the use of a varied binary combinations of GeI2, GeCl2, Ge(HMDS)2, GeI4 and Sn(HMDS)2, SnCl2, (Bu)2SnCl2, (Bu)4Sn and SnI2 for tin. To synthesize SnxGe1-x NCs in any appreciable amount, we found it necessary to either combine a highly reactive Ge precursor and relatively inert Sn precursor, or a pair of precursors that in-situ form a complex which eventually can be reduced and/or thermolyzed to SnxGe1-x alloy NCs at elevated temperature in the presence of a reducing agent. Of these, the combination of GeI2 and Sn(HMDS)2 with n-BuLi was found to offer the best combination of crystallinity and control over size and composition in the product SnxGe1-x NCs. (Scheme 1). As a pale yellow GeI2/Sn(HMDS)2 solution in OLA is heated from room temperature to 100oC, the solution color changes to a bright orange, which subsequently turns reddish-brown and black as temperature upon further heating to 180oC. Importantly, this pattern of color change is not produced by similar heating of either precursor in OLA alone (Figure S1). Such accelerated darkening of the reaction mixture is believed to indicate nucleation of very small alloy parti-


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cles, likely through disproportionation of the Sn(II) and Ge(II) precursors into corresponding (IV) and (0) species in the presence of amine, as was observed in previous syntheses of pure Ge NCs from Ge(II) precursors.29 SAXS measurements do not reveal the presence of any scattering entities in the reaction mixture up to 180oC, indicating a lack of particles above ~1nm. However, continued heating to a targeted growth temperature ≥ 200 oC eventually produces NCs in very low chemical yield with Sn content limited to ≤18% without any addition of a reductant. We refer to this approach as the self-reduction method (Figure S2). Addition of n-BuLi on an equimolar basis with the Sn-precursor at the target temperature was found to greatly enhance both the chemical yield and the range of attainable Sn concentrations in the product NCs. Varying the ratio of Sn-to-Ge precursors from 0 (i.e., only Ge) to 1 offers control over the Sn-content from 0-42% Sn (Figure S3). Reactions with higher Sn:Ge ratios than 1:1 yielded NCs with somewhat lower (only up to ~ 22 %) Sn incorporation along with the formation of -Sn impurities. This pattern of color changes and product formation suggests that NC growth proceeds via formation of heterometallic complex(es) between the original Sn and Ge precursors, which is followed by the (self-) reduction of these intermediate entities at elevated temperatures to form Ge-Sn nuclei, which subsequently proceed to grow through typical colloidal processes. Evidence of complex formation includes the diminishment of a strong singlet at 678ppm in the 119Sn NMR of a Sn(HMDS)2 solution in OLA (Figure 1), which was previously attributed to the Sn-amine complex,30 upon gradual addition of GeI2/OLA solution. This is accompanied by the appearance of a new signal at 300ppm, upfield from the initial singlet. As more GeI2 is added and the Sn:Ge ratio reaches approximately 2:1, the resonance at 678 ppm disappears completely and only the signal at 300ppm is evident (Figure 1). As the Sn:Ge ratio is further decreased to 1:1, the resonance at 300ppm also vanishes. In similar NMR measurements with the reverse system, SnI2/Ge(HMDS)2 in OLA, the peak at -380 ppm for pure SnI2/OLA is not appreciably affected by the addition of Ge(HMDS)2. Further study of the more detailed reaction pathway is underway, but these preliminary results suggest the formation mixed-ligand homometallic or heterometallic complex(es) at room temperature before the formation of the Ge-Sn nuclei occurs at elevated temperatures (Scheme 1). It is worth noting that the use of oleylamine is crucial for the successful synthesis of alloy nanocrystals. Other solvents/capping agents such as trioctylphosphine, octadecene or even a mixture of oleylamine (1 mL) and triocylphosphine (4 mL) or octadecene (4



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mL) did not yield the desired product. Finally, use of the nominally hydride-donating reducing agent “Super-hydride” yielded insoluble products with low concentration of Sn suggesting hydride ion inhibits the growth of NCs, perhaps by forming Ge-H termination on the surface of SnxGe1-x nanocrystals. Overall, we find that the variations between our approach and that applied by Arachchige28 seem to result in a less complex reaction that offers higher chemical yields and greater flexibility for producing SnxGe1-x NCs with narrower size dispersity over a wider continuous range of sizes and Sn concentrations (up to 42%). Structural and Morphological Investigations of SnxGe1-x Alloy Nanocrystals. The shape, size and crystallinity of synthesized NCs were analyzed using a combination of TEM, XRD and SAXS. NCs are on average quasi-spherical (Figures 2a-d) with average diameter ranging from 5 to 15 nm, depending on the reaction temperature and the duration of particle growth. The average size of NCs reported throughout this paper is based on SAXS data. For several samples SAXS-based size was corroborated with measurements from TEM images, showing very good agreement in both size and size dispersity (Figure 3). According to XRD analysis, the crystal structure of SnxGe1-x NCs is still cubic even at high concentrations of Sn (Figure 4). The major diffraction peaks are indexed as (111), (220) and (311) planes of diamond cubic structure with Fd m space group (ICDD: 04-0545). The measurement confirms that the NCs are free from tetragonal Sn ( -Sn), which is the form typically seen in the bulk at room temperature and in previously reported pure Sn NCs,30-31 and also free of crystalline oxides of Ge and Sn. The lattice constant values in SnxGe1-x nanocrystals vary from 5.646 Å for x = 0 to 5.958 Å for x = 0.42. The concentration of Sn in NCs, x, is extracted from the lattice constant using a quadratic relation a(SnxGe1-x)=a(Sn)x+a(Ge)(1-x)-bx(1-x) and bowing parameter b = 0.166Å established previously.32 The Sn-content was also independently verified using elemental analysis by EDX and XPS (Table 1). High-resolution TEM images reveal (Figure 2e-f) predominantly singe-crystalline particles, which is consistent with the size obtained from Scherer line-broadening analysis in XRD and SAXS. Lattice spacing in the (111) direction determined using intensity line scans across the inverse FFT of high-resolution images (Figures 2b) were in good agreement with those observed in XRD measurements. Although the solubility of Sn in bulk cubic Ge is quite low (~1%),33 our observations are consistent with the persistence of a highly crystalline, homogenous SnxGe1-x alloy even at Sn


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concentrations as high as 42%. To further exclude significant segregation of Sn and Ge into separate domains, NCs were imaged with HAADF microscopy, a “Z-contrast” technique providing information on elemental distribution based on atomic number (Figures 5a-b), in combination with STEM-EDX mapping. Figure 5c-e depicts overlapped areas of Ge- and Sn-selective elemental mapping over the area indicated by a box in Figure 5b, showing a homogeneous distribution of either element over the entirety of each NC. The existence of a single-component distribution of Sn and Ge in the NCs was established by applying a multivariate statistical analysis34 on the STEM-EDX mapping images. The survey XPS spectrum exhibits peaks corresponding to Ge, and Sn, along with C with no other impurities present, which is consistent with EDX survey spectrum as well. Typical high resolution XPS spectra of Ge 3d, and Sn 3d5/2 electrons from 9 nm sample of Sn0.27Ge0.73 alloy NCs confirm the expected distribution of Sn and Ge elements (Figure 6). The Ge 3d peak was fit with three Gaussian/Lorentzian components at 28.64 ± 0.01 eV, 29.62 ± 0.01 eV, and 31.20 ± 0.01 eV with FWHM of 1.130 ± 0.003 eV, 2.0 eV, and 1.3 eV, respectively, that can accordingly be assigned to Ge(0) of alloy, Ge(II), and Ge(IV) species. The Sn 3d5/2 peak was fit with two Gaussian/Lorentzian components at 484.22 ± 0.08 eV and 485.45 ± 0.12 eV (FWHM of 0.85 ± 0.05 eV and 1.9 eV), respectively. The first narrow component can be assigned to Sn(0) of alloy, whereas the broad second peak is attributed to Sn(II)/Sn(IV) species, likely from ligand-coordinating surface atoms. In these XPS measurements, the average Sn/Ge ratio was found to be 0.43, which corresponds to x=0.30 and is in good agreement with the x=0.27 that was determined from XRD measurements. Ge and Sn in higher oxidations states (II, and IV in particular) can also be indicative of post-synthesis oxidation, and indeed a relatively high overall fraction of XPS signal from these species has been directly related to formation of oxides, most probably at the surface of alloy NCs.28 The large relative contribution from Ge(0) and Sn(0) in our measurements suggest very little, if any, oxide formation, which is consistent with our high-resolution TEM/EDX data. Temperature- and Time-Dependent Studies on formation of SnxGe1-x Alloy Nanocrystals. For the overall goal of achieving the maximum concentration of Sn in crystalline SnxGe1-x NCs, there also appears to be an ideal reaction temperature of 200oC. Below this temperature, a large fraction of the product was found to be amorphous. If higher growth temperatures are used, larger NCs are obtained, but the Sn-content invariably decreases, all other reaction conditions being held the same. For instance, an increase in growth temperature from 200oC to 280oC in-



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creases NC size from 4 nm to 10 nm, but reduces the concentration of Sn from 40-42% to below 20%, as determined after 30 min of particle growth (Figure 7; Table 2). Intriguingly, a similar but much less dramatic trend was observed with an increase in growth time at a given temperature. NCs grown for 120 min at 230oC show a consistent 1-2 % depletion of Sn relative to those grown for 15 min, even though the particle growth is significant only over the first 20-30 min of the reaction (from 6 nm to 9 nm, Figure S4). At the same time, prolonged growth times past the period of active particle growth (longer than 30min) lead to narrower size distribution of final nanoparticles. For example, annealing the product for 2 hours at 230oC reduced its size distribution from 40-45% to 20-23% according to SAXS data (Figure S5. Table S1). Annealing the product beyond 2 hours does not lead to further improvements in size distribution, and eventually leads to its broadening through Ostwald ripening. Thermal Stability of SnxGe1-x Alloy Nanocrystals. Temperature and time-dependent studies of nanocrystal growth suggest that even at the growth stage there is a net diffusion of Sn out of the particle interior upon prolonged exposure of alloy NCs to temperatures above 200oC. Indeed, the synthesis of NCs at 280 °C leads to the formation of larger (11-12 nm) predominantly single crystalline particles of SnxGe1-x with reduced Sn concentration and the discernible presence of elemental -Sn phases (according to X-ray diffraction in Figure 7). To further understand the diffusion out process of tin from the alloy lattice, we have annealed the isolated and purified alloy NCs nanocrystals under nitrogen at high temperatures under thermogravimetric analysis (TGA) conditions [Figure 8(a)]. XRD patterns of samples annealed at 200 °C show a small shift in peak positions towards higher 2 angles, indicating the lattice shrinking due to the some loss of Sn. The loss accelerates at higher temperatures, and eventually reflections due to the -Sn phase emerge, although samples annealed at 500 °C still retained ~5 % Sn in the cubic lattice (in this case, from an initial 18 %). Annealing at 900oC for 30 min, however, leads to the complete separation of two elements [Figure 8(b)]. Changes in the NC lattice observed by TEM imaging as the sample was heated are consistent with the conclusions above. As sample temperature increased, the (111) distance in the alloy lattice gradually decreased until it reached a value indistinguishable from that of pure Ge at 500 oC (Figure 9). As evident from Figure 9, even at 500oC no sintering of NCs was observed and they clearly retained their separation and morphology. Lattice instability and the high mobility of Sn


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atoms above the Sn melting point are likely the factors causing the observed phase segregation. It is possible that controlled oxidation of the NC surface could be used to slow down or halt the Sn diffusion, but adverse and uncontrolled effects oxidation can have on optical and/or electronic properties of the final product render this a far from ideal approach. It is important to note that this interpretation stands in contrast to that advanced in a previous report on similar materials.28 Arachchige and coworkers attributed a very similar shifting of diffraction peaks to higher angles at temperatures up to 500 °C to a thermally-induced transition from a tensile-strained to a compressively strained lattice, without any change in alloy composition or phase segregation. Based on our experience with other NC materials, in a free-standing pseudo-spherical NC with essentially homogeneous composition, in the absence of an external structure-directing influence (e.g., a substrate in epitaxial contact), such a transition seems unlikely. Thus, particularly given the corroborating observations of a decline in Sn content with prolonged reaction time and distinct signatures of -Sn formation at 500 °C, we find migration and loss of Sn a more satisfying interpretation of the XRD data. Microwave Synthesis of SnxGe1-x Alloy Nanocrystals. Given the complex effect of temperature on size and composition of these SnxGe1-x NCs, we also investigated the use of microwave instead of convection heating, inspired by the recent observation that microwave approach favors the synthesis of highly crystalline Ge NCs from GeI2.35 Indeed, we found that microwave heating can also be used to drive the synthesis of SnxGe1-x alloy NCs. XRD patterns indicate a very similar range of attainable Sn-contents (Figure S6), and similar average sizes were observed. However, in our experiments, microwavesynthesized NCs tended to be more irregularly shaped, and to have wider size distributions (Figure 10 and S7). This may indicate that the above-mentioned annealing processes that lead to sizefocusing may, as would be expected, also preferentially lead to more spherical shape, and that preferential absorption of microwaves by the precursors relative to the product NCs means suppression of this annealing. Optical Properties of SnxGe1-x Alloy Nanocrystals. Figure 11a shows Vis-IR linear absorption spectra for typical alloy NCs of the same size (9 nm (d)) suspended in CCl4 over the wavelength range of 500 nm to 2500nm. Although they appear featureless across the whole range of wavelengths, similar to pure Ge nanocrystals,19 we



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observe the anticipated bathochromic shift with increasing Sn-content, as previously reported for SnxGe1-x thin films.24, 36 Through further analysis of these spectra, we attempt to extract the effective optical band gap (Eg) for these materials and to look for evidence of an indirect-to-direct transition. We first applied the method commonly used for non-crystalline semiconductors offered by Tauc in his work with H-passivated amorphous silicon.37 Relying on the assumptions of parabolic conduction and valence band extrema and a constant value of the momentum matrix element during the optical transition, Tauc demonstrated that the extrapolation of linear portions in plots of (Aħ )2 and (Aħ )½ vs ħ to the energy axis would allow one to extract values for the direct and indirect band gaps, respectively. Attempts to establish values for Eg for direct transitions via this method led to ambiguous results, as the identification of a linear regime in the (Aħ )2 vs ħ plot clearly depended on the plot magnification scale (Figure S9). Indeed, although a previous report used the Tauc technique to determine values for direct band gap for alloy particles of a more limited range of Sn-content (up to 12%), they noted appreciable variations from a monotonic red-shift.28 While these variations were ascribed to unaccounted-for effects of band structure changes and quantum confinement, we hold that uncertainty in the identification of linear regimes within the data provided therein could be an equally important factor. Finally, we note that the absorption spectra in Ref. 28, which were extracted from diffuse reflectance measurements, have noticeably different shapes than our directly measured absorption spectra, as well as those reported elsewhere for pure Ge NCs38. Because its correct application requires reliable data over a significant spectral region, a Tauc analysis of converted reflectance data may produce quite different results, further diminishing confidence in band gap values derived in this manner. In Figure 11b, we see that assignment of an indirect band gap is marginally more successful, and produces the expected trend in Eg vs. Sn-content, although the linear regimes are fairly limited in extent. Overall, we find that analysis of our spectra by the method of Tauc does not allow us to definitively state whether any of our alloy NCs are clearly direct or indirect. It is likely that a combination of factors, including the potential contribution of intra-gap surface states, and spectral interference from ligands are obfuscating the spectra and contributing to this ambiguity. Setting aside the question of indirect vs. direct, we note that Cody proposed a further modification of the Tauc model for the same material (amorphous silicon), replacing the momentum conservation assumption with maintaining a constant dipole matrix element.39 In this case, a


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value for optical band gap value can be derived from similar extrapolation of the linear regime within a plot of (A/ħ )½ vs ħ . As seen in Figure 11c, linear behavior in such plots is observed over a wider range of energies, allowing a somewhat more confident assertion of Eg for each NC composition that confirms the same trend observed in the Tauc plots. Our Eg values are somewhat higher than those obtained in SnxGe1-x thin films of similar composition, 24, 36 which is consistent with the expected influence of quantum confinement effects. Similar analysis has been conducted with samples of different size but the same composition (Figure 12). Linear absorption, Tauc and Cody plots are presented for nanoparticles of Sn0.10Ge0.90 with sizes ranging from 7.7 to 11.6nm, demonstrating the expected size-dependent shift in all three sets of plots. It is worth reflecting on that while the size-dependence of the optical absorption of alloy NCs of relatively low Sn-content are in accord with expectations of the quantum confinement effect, our attempts to observe confinement-assisted direct gap behavior in SnxGe1-x NCs of higher Sn-content have yielded less conclusive results. This could potentially be caused by complications arising from the distinctive characteristics of the electronic structure of these materials at the nanoscale. The offset between the two lowest conduction band minima (L and ) in Ge is ~0.15 eV. In Sn, the

minimum is very slightly lower, which is the reason Sn alloying may be

expected to produce a transition to direct-gap behavior in a bulk material.26 Less predictable, however, is the effect of quantum confinement on the relative positions of these two band minima in SnxGe1-x NCs. In each pure material, the effective mass of electrons (me) at the

point is

lower than that at the L point, suggesting that the same may be true in SnxGe1-x alloys. This raises the possibility that depending on the exact composition, confinement effects may actually inhibit a crossover to direct-gap behavior. Moreover, as the concentration of Sn in the alloy increases, the hybridization of Ge and Sn bands is likely to yet again alter the qualitative band diagram of the alloy relative to the pure band structures. As such, further investigation of the alloy system is needed to fully understand the influence of size on the indirect-to-direct band gap transition. In addition, several samples with high Sn-content exhibited broad, weak photoluminescence at energies near the optically determined band gap (Figure 13a), a phenomenon that has previously been taken as a sign of at least partial direct-gap character of the principal transition in other Group IV NCs26. We must allow that it is also possible that the emission originates from the surface trap states of the nanocrystals and further experiments, particularly in improving sur-



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face passivation (such as in the manner recently applied to pure Ge NCs),40 are required to reveal the true origin of the photoluminescence together with understanding the set of factors limiting emission efficiency. Given that the primary motivation for investigating a potential indirect-todirect transition is to arrive at a material with enhanced light absorption, we investigated the Sncontent dependence of per-NC molar extinction coefficient, , using previously reported procedure.41 Figure 13b presents

as a function of photon energy for SnxGe1-x NCs with x=0.00 (i.e.,

pure Ge), 0.10, 0.27, 0.36 and 0.40. A very clear and significant increase in

across the entire

spectrum is evident as the concentration of Sn increases, an effect that is particularly pronounced in the solar-relevant region between 1-1.5 eV. At 3.1 eV,

for NCs of high Sn-content reaches

values comparable to those measured in other direct-gap NCs of comparable Eg, such as PbSe42 and Ag2Se.41 █Conclusion:

In conclusion, we have developed the first facile synthetic approach to colloidal SnxGe1-x homogeneous alloy NCs with Sn concentrations reaching as high as 42%, and sizes ranging from 5-15 nm with narrow size distribution. We have found that no reduction agents are needed in the synthesis of SnxGe1-x alloy nanocrystals with x below 0.18, although the yield of the final product is low. The use of n-BuLi as the reducing agents can increase in tin content up to 42% in the final product. The use of microwaves in the synthesis of the alloy NCs did not show any particular benefits over the convection approach. In fact, due to the lower temperatures accessible via microwave heating and lower quality of the final product, we advocate to use the convection heating for the synthesis of the allow NCs.

119

Sn NMR measurements showed that the possible

formation of complex(es) involving GeI2/Sn(HMDS)2/OLA. The systematic shift in X-ray diffraction peak positions on increasing tin content, HAADF and EDX elemental mapping analysis confirmed the uniform distribution of elements SnxGe1-x alloy lattice. The NCs maintain a diamond-like cubic structure, but appear to be metastable, undergoing Sn-depletion and then eventually full phase segregation as temperatures increase from ~220-900oC. Absorption spectroscopy of NCs of varying composition demonstrates a pronounced red-shift and increase in molar absorptivity as a true indicator of solar relevance with increasing Sn-content. Although definitive evidence of a transition to direct-gap behavior remains elusive, weak photoluminescence is observed in a substantial fraction of NC samples in which Sn-content exceeds 35%. These results


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suggest that solution-processible NCs based on alloys of non-toxic and abundant Group IV elements are a potentially promising material system for solar energy capture applications.

Supporting information available: Additional TEM images, XRD and SAXS data and Absorption spectra. This material is available free of charge via the Internet at http://pubs.acs.org Corresponding Author *[email protected] Acknowledgement J.M.P. acknowledges the support of the Center for Advanced Solar Photophysics (CASP), an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences. A.F.F. is supported by a LANL Director’s Postdoctoral Fellowship. K. R. is supported by a LANL LDRD funding. This work was performed, in part, at the Center for Integrated Nanotechnologies, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science. Los Alamos National Laboratory, an affirmative action equal opportunity employer, is operated by Los Alamos National Security, LLC, for the National Nuclear Security Administration of the U.S. Department of Energy under contract DE-AC52-06NA25396. Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin Company, for the US Department of Energy under contract DEAC0494AL85000.

_____________

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