Letter pubs.acs.org/NanoLett
Core-Shell Germanium/Germanium−Tin Nanowires Exhibiting RoomTemperature Direct- and Indirect-Gap Photoluminescence Andrew C. Meng,† Colleen S. Fenrich,† Michael R. Braun,† James P. McVittie,‡ Ann F. Marshall,†,§ James S. Harris,∥ and Paul C. McIntyre*,† †
Department of Materials Science and Engineering, ‡Stanford Nanofabrication Facility, §Stanford Nano Shared Facilities, and Department of Electrical Engineering, Stanford University, Stanford, California 94305, United States
∥
S Supporting Information *
ABSTRACT: Germanium−tin alloy nanowires hold promise as siliconcompatible optoelectronic elements with the potential to achieve a direct band gap transition required for efficient light emission. In contrast to Ge1−xSnx epitaxial thin films, free-standing nanowires deposited on misfitting germanium or silicon substrates can avoid compressive, elastic strains that inhibit formation of a direct gap. We demonstrate strong room temperature photoluminescence, consistent with band edge emission from both Ge core nanowires, elastically strained in tension, and the almost unstrained Ge1−xSnx shells grown around them. Low-temperature chemical vapor deposition of these core-shell structures was achieved using standard precursors, resulting in Sn incorporation that significantly exceeds the bulk solubility limit in germanium. KEYWORDS: Germanium−tin, core−shell nanowire, photoluminescence, optoelectronics
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nanowire structures can accommodate large elastic strains through dilatation.17 Surface quality is an often overlooked property for optoelectronic devices and is an important issue for Ge1−xSnx nanowires. Given the low equilibrium solubility of Sn in Ge and the low surface energy of Sn, there will be a driving force for Sn segregation to the wire surface and this may cause Sn precipitates to form, leading to nonuniform nanowire composition. Previous work has demonstrated photoluminescence (PL) from Ge1−xSnx nanowires at a temperature of 7 K and from epitaxial two-dimensional (2D) Ge1−xSnx thin films under various conditions2,6,18 but has not focused on possible surface effects or on the detailed features of the spectra. For pure germanium, it is known that nanowires with only a native oxide coating do not exhibit indirect-gap PL due to surface recombination19 while core-shell Ge/SiGe nanowires exhibit more bulk-like PL, including both direct- and indirect-gap emission because confinement in the Ge core effectively separates photogenerated carriers from surface defects.20 It is not known whether similar effects occur for Ge1−x Sn x nanowires. Ge/Ge1−xSnx core-shell nanowires were synthesized using Au-catalyzed vapor−liquid−solid growth (Figure 1), achieving average Sn compositions up to 7 at. % using germane (GeH4) and tin(IV) chloride (SnCl4) precursors in hydrogen carrier
e 1−x Sn x with high Sn composition has attracted considerable interest as a potentially silicon-compatible direct band gap semiconductor for integration of electronic and photonic devices.1−7 As transistors continue to scale to smaller dimensions, energy dissipation in complementary metal oxide semiconductor (CMOS) circuits is increasingly dominated by electronic interconnects, a problem that may be alleviated by transitioning to optical interconnects.8−10 Although III−V hybrid devices offer another approach,11 there are implementation barriers due to incompatibility of III−V growth processes with underlying silicon circuitry and the difficulty of preparing very high crystalline quality III−V epitaxial thin films on lattice mismatched substrates, such as group IV semiconductor crystals.2,12 Although Ge1−xSnx offers advantages in these regards, there are two significant challenges to growth: first, the indirect-to-direct band gap transition occurs at compositions in the range ∼6.5−11 at. % Sn for unstrained Ge1−xSnx,2,13 which is significantly higher than the thermodynamic solubility limit of Sn in Ge diamond cubic crystals; second, α-Sn has a 14% larger lattice parameter than Ge, which can lead either to formation of misfit and threading dislocations12 that cause nonradiative recombination losses or, if coherent interfaces are achieved in Ge1−xSnx nanostructures, to retained compressive elastic strain that suppresses formation of a direct band gap.13,14 In principle, vapor−liquid−solid growth of Ge1−xSnx nanowires can address these challenges because it is a low-temperature, catalyzed growth process,15,16 which facilitates supersaturation of Sn, and free-standing © XXXX American Chemical Society
Received: August 7, 2016 Revised: October 23, 2016 Published: November 1, 2016 A
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(Figure 2b); the diameter at the base is approximately 180 nm, and the maximum diameter is approximately 270 nm (Figure S1 in Supporting Information). We observe that the solidified catalyst particle is larger after Ge1−xSnx shell deposition, consistent with the presence of residual Sn in the particle after the cessation of growth. The increase in wire length after Ge1−xSnx (stage 3) growth is modest and not significantly greater than the radial thickness of the shell, despite that an identical 30 min growth duration resulted in deposition of the initial 4.0 μm long Ge wires. The SEM images of nanowires grown on Ge(111) suggest that the distance between nanowire tip and its surrounding substrate surface region decreases as a function of Sn precursor partial pressure during shell growth from 4.0 ± 0.2 μm at PSnCl4/PGeH4 = 0 to 3.8 ± 0.2 μm at PSnCl4/ PGeH4 = 2 × 10−5 (Figure 2). This does not reflect an actual decrease in the nanowire length but rather (1) the fact that there is some uncertainty in the length measurement related to a possible angle offset for different samples mounted in the SEM and (2) that there is an increase in the rate of epitaxial chemical vapor deposition (CVD) on the surrounding substrate due to catalytic effects from the presence of gaseous Sn species.21 The pure Ge nanowires grow primarily in the vertical [111] direction, parallel to the surface normal of either Ge or Si single crystal substrates, providing the same epitaxial orientation relationship to the Ge/Ge1−xSnx core-shell nanowires. On Si(111) substrates, approximately 60% of the core-shell Ge/ Ge1−xSnx nanowires grow in inclined ⟨111⟩ directions (Figure 2c). Hexagonal faceting is observed on both the Ge nanowires and along most of the length of the core-shell nanowires. The facets are accentuated on the Ge1−xSnx shell sidewalls. Strain Characterization by X-ray Diffraction. XRD symmetric (ω-2θ, ω = θ) scans of as-grown samples of Ge nanowire reference samples and Ge/Ge1−xSnx core-shell nanowires of different nominal Sn composition deposited on Si(111) substrates show the substrate (111), (222), and (333) peaks as well as (111) and (333) peaks for epitaxial nanowires (Figure 3a). The (222) peak is a diffraction artifact resulting
Figure 1. Core-shell Ge/Ge1−xSnx nanowire growth schematic showing different stages of growth.
gas. Room-temperature PL measurements indicate strong emission signatures matching direct- and indirect-gap transitions in both the tensile strained Ge core and the almostunstrained Ge1−xSnx shell. These emission features occur despite the presence of defects on the shell surface and nonuniform incorporation of Sn in the single crystal shell. The strong room-temperature PL is a promising sign for optoelectronic applications of Ge1−xSnx nanowires. Assignment of PL features takes into consideration both the measured Sn composition of the shell, and core and shell lattice constants measured by X-ray diffraction (XRD). Comparison of Ge Nanowires and Ge1-xSnx Core− Shell Nanowires. SEM images of vertical Ge nanowires obtained by two-stage growth show minimal taper after the initial, higher-temperature, seeding stage of growth and had diameters of ∼50 nm (Figure 2a, Figure S1 in Supporting Information). In contrast, Ge/Ge1−xSnx core-shell nanowires obtained by three-stage growth (Figure 1) exhibit taper at the tip and inverse taper along the majority of the wire length
Figure 2. Low-magnification SEM and magnified inset of (a) control Ge nanowires, (b) core-shell Ge/Ge1−xSnx nanowires PSnCl4/PGeH4 = 2 × 10−5 grown on Ge(111) and (c) Si(111). Inset scale bars are 1 μm; primary scale bars are 10 μm. B
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Figure 3. (a) Wide-angle XRD of as-grown epitaxial Ge/Ge1−xSnx core-shell nanowires on Si(111) substrate. (b) XRD, off-axis scan, highly vertically aligned as-grown epitaxial Ge/Ge1−xSnx core-shell nanowires on Ge(111) substrate. The solid and dashed lines indicate the calculated peak positions for a sample with 4 at. % Sn according to the elasticity model and fully relaxed model, respectively, as a reference mark.
coherency strain is stored in the core for a core-shell nanowire with a core radius of 25 nm and shell thickness of 75 nm. Because of the minimal compressive strain in the Ge1−xSnx shell, estimated to be less than −0.092% from elasticity calculations, these data cannot differentiate between fully relaxed regions and strained shell regions. This is shown by the solid and dashed lines in Figure 3b, indicating the calculated peaks for the coherent and fully relaxed models, respectively, would lay almost on top of one another. The result of this near overlap is that any scattering from relaxed Ge1−xSnx deposition on the (111) substrate surface between the nanowires is convoluted with scattered intensity from the Ge1−xSnx shells. However, this leads to an uncertainty of only ∼0.2 at. % in the average Sn concentration determined by X-ray diffraction. Furthermore, we observe that the average composition inferred for the deposited Ge1−xSnx saturates at approximately 3.5−4 at. % for the CVD growth conditions employed. These data can then be compared to composition measurements made by X-ray energy dispersive spectroscopy (EDS) in TEM, and also used to analyze the effects of strain and composition on the Ge1-xSnx band structure as probed by photoluminescence. Transmission Electron Microscopy (Local EDS Sn Composition Measurements, Crystallographic Quality of Ge1−xSnx Shells, and Defect Characterization). Lowmagnification TEM images with corresponding convergent beam TEM-EDS composition measurements (Figure 4) were taken for Ge1−xSnx nanowires grown with different SnCl4 partial pressures. GeS2 and SnS calibration standards were used to obtain corrections for the Cliff−Lorimer k-factors for Ge and Sn in TEM-EDS (Figure S3, Table SI in Supporting Information).24 For nanowires grown at PSnCl4/PGeH4 = 1 × 10−5, 2 × 10−5, and 3 × 10−5, the average Sn compositions along the wires were 3 at. %, 6 at. %, and 7 at. %, respectively (for spectra, see Figures S4−S6 in Supporting Information). Although the data indicate a slightly higher Sn composition near the tip of the Ge1−xSnx nanowires than at the base, the composition variation along the wire length is largely within the measurement uncertainty. The data suggest a high degree of Sn incorporation into the core-shell Ge/Ge1−xSnx nanowires above the equilibrium solubility limit of Sn in Ge (∼1 at. %); however, these results are averaged over volumes of 0.001−0.008 μm3 at the positions at which the EDS analysis is performed. The Sn composition measured by EDS is significantly higher than that
from symmetry breaking at the wafer surface. No additional peaks were observed. “Off-axis XRD” (ω-2θ, ω = θ + Δω with a misalignment of Δω = 0.1°) was performed on both Ge nanowire and Ge/Ge1−xSnx core-shell nanowire array samples grown on Ge(111) substrates (Figure 3b). The slight off-axis tilt greatly attenuates the intense substrate peaks with minimal change in nanowire peak intensities. Greater than 95% of nanowires on these samples were vertically aligned, thus minimizing complications from azimuthal and radial strain broadening due to inclined nanowires.22 Axial strains are detected by measuring lattice spacings normal to the substrate surface. Furthermore, we assume the interface between the core and shell is coherent if there are no misfit dislocations detected (see Figure 5). In the fully-strained case, the nanowire peak in Figure 3b corresponds to combined X-ray scattering of both the core (∼25 nm radius) and shell regions (∼75 nm thickness). Tin compositions can be inferred from the data using a continuum elasticity model (see Figure S2 in Supporting Information),23 which is referred to here as the “coherent model”. On the other hand, if the core and shell were fully strain relaxed by misfit dislocations, the nanowire peak in Figure 3b would correspond only to X-ray scattering from the Ge1−xSnx nanowire shell. Although this is not the case for the samples studied, an assumption of such a relaxed state can also be applied to analysis of the diffraction results. Therefore, in the limits of either full core-shell misfit strain relaxation or of coaxial elastic strains resulting from fully coherent core-shell interfaces,23 it is possible to estimate the bulk Sn composition in the GexSn1−x shell based on the XRD peak positions and Vegard’s law. Nanowires grown with PSnCl4/PGeH4 = 1 × 10−5, 2 × 10−5, and 3 × 10−5 correspond to Sn compositions of 2.9 at. %, 3.7 at. %, and 3.5 at. %, respectively, under fully relaxed conditions (Table 1); the coherent model assumptions give Sn compositions of 3.1 at. %, 3.9 at. %, and 3.7 at. %, respectively. These values are similar despite very different assumptions about the coherency of core and shell because almost all of the Table 1. Sn Composition in Ge1−xSnx Shell PSnCl4 /PGeH4
fully relaxed (at. % Sn)
coherent (at. % Sn)
1 × 10−5 2 × 10−5 3 × 10−5
2.9 3.7 3.5
3.1 3.9 3.7 C
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Figure 4. Low-magnification TEM images of Ge/Ge1−xSnx core-shell nanowires and EDS measurement points for (a) PSnCl4/PGeH4 = 1 × 10−5, (b) PSnCl4/PGeH4 = 2 × 10−5, and (c) PSnCl4/PGeH4 = 3 × 10−5. All scale bars are 1 μm. (d) Corresponding EDS composition data; the remaining composition at position 5 is Au VLS catalyst.
planes in the shell are shown in Figure 7. Although there are planar contrast features visible in cross-section in the low magnification TEM image in Figure 7a, HRTEM (e.g., Figure 7b) shows that the lattice fringes are continuous across these planar defects, indicating that they form coherent interfaces with the adjacent non-defective Ge1−xSnx crystal. Because the regions over which EDS measurements are taken are 100−200 nm in both lateral size and depth, the compositions detected by EDS, which average non-defective areas with defective ones, are expected to be higher than compositions inferred from homogeneous strains measured by XRD. Near one particular coherent (111) planar defect, local EDS measurements using convergent beam shows an increase in local Sn composition to ∼19 at. %. The Sn-rich coherent defects are not observed in XRD. Consistent with the TEM analysis showing a coherent interface with the surrounding Ge1−xSnx, this suggests that these defects have the same crystallographic registry25−28 as the surrounding material; the defects occupy too small a volume fraction compared to the rest of the epitaxial Ge/Ge1−xSnx core-shell nanowires to be detected using lab-source XRD. Sn-rich precipitates observed on the surface of the nanowires in TEM also are not detected in XRD. These precipitates appear to be randomly oriented and constitute an even smaller volume fraction than the coherent defects. Photoluminescence Spectrum Deconvolution and Excitation Intensity Dependence. Room-temperature pump power dependent PL spectra of as-deposited core-shell Ge/Ge1−xSnx nanowires grown on Si(111) substrates at PSnCl4/ PGeH4 = 2 × 10−5 (4 at. % in the shell as inferred from XRD measurement) exhibited emission features at four different photon energies fit to the following equation where I is the fitted PL spectrum intensity, I0 is a constant offset, and Ai and
predicted by XRD. This is consistent with incorporation of Snrich phases in the nanowires that are not detected by XRD. In fact, Sn-rich precipitates approximately 10−20 nm in diameter are observed at the surface of the Ge1−xSnx shells and Sn-rich coherent planar defects are observed within the shells (vide infra). These defects do not appear to affect the crystalline quality of the Ge1−xSnx matrix either beneath the surface precipitates or surrounding the planar defects. Phase contrast high-resolution TEM (HRTEM) images of the shell region of Ge/Ge1−xSnx core-shell nanowires taken with the beam direction along the [112]̅ zone axis are shown in Figure 5. The (111) lattice fringes perpendicular to the growth axis of the wires and (2̅20) lattice fringes are clearly visible for all of the nanowires imaged, and the corresponding diffraction spots are labeled in the fast Fourier transforms of the images (Figure 5). These images and all others acquired from the samples show continuous lattice fringes for the Ge1−xSnx shells and no evidence of stacking faults or core-shell misfit dislocations. The nanowires are single crystals. Combined with the SEM images, the TEM results indicate that high Sn composition Ge1−xSnx nanowire structures with a Ge/Ge1−xSnx core-shell structure were successfully grown; however, the effective bulk incorporation of Sn in the shells appears to saturate at ∼4 at. % based on the XRD analysis. We observe two types of defects in core-shell Ge/Ge1−xSnx nanowires. Under all Sn precursor partial pressure flows, small Sn-rich precipitates and Sn-rich planar coherent defects are seen in the Ge1−xSnx shell region (Figures 6 and 7). Similar coherent Sn defects have previously been reported in both Si and Ge crystals. 25−28 The small precipitates can be distinguished by the Moiré fringes in the TEM image, and EDS shows Sn compositions greater than 50 at. % in the precipitates (Figure 6a, b). Coherent Sn-rich segments of (111) D
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Figure 5. Ge1−xSnx shell [112̅] zone axis high resolution TEM images with FFT inset (a) PSnCl4/PGeH4 = 0, (b) PSnCl4/PGeH4 = 1 × 10−5, (c) PSnCl4/ PGeH4 = 2 × 10−5, and (d) PSnCl4/PGeH4 = 3 × 10−5. Images are taken near the tip of the nanowires. All scale bars are 5 nm.
with the adventitious layer did not rise above the noise level. Lack of signal from these areas suggests that the nanowires are responsible for the PL. No changes in the sample were observed in SEM after collection of PL spectra. A micro-PL setup with an 8 μm diameter spot was used. The highest pump fluence of 60.0 kW cm−2 corresponded to only a 30 mW power at the sample surface. Least squares fitting was used to locate the emission peak centers, shown in Table SII in the Supporting Information. Fit parameters are shown in Tables SII and SIII in the Supporting Information. For these Ge/Ge0.96Sn0.04 core-shell nanowires at 60.0 kW cm−2 laser pump power intensity, the peak centers are 0.550, 0.578, 0.631, and 0.695 eV with a reduced χ-squared value of 19.9 for the fit (Figure 8a, see also Table SII in Supporting Information). These correspond to emission wavelengths of 2255, 2145, 1965, and 1784 nm. Fitting the PL data to four peaks not only best reproduces the spectral line shape, which comprises two central peaks and two broad shoulders at the lower and higher wavelength extremes, but also results in the largest decrease in reduced χ2 values with respect to the number of additional peaks being fit. The four PL emission features in Ge1−xSnx nanowires also have physical significance and can be indexed to the Ge1−xSnx Lvalley (k = ⟨111⟩, indirect gap transition), Ge1−xSnx Γ-valley (k = ⟨000⟩, direct gap transition), Ge L-valley, and Ge Γ-valley transitions in order of increasing photon energy. First, we consider the spectrum from pure Ge nanowires (PSnCl4/PGeH4 =
Figure 6. Sn-rich precipitates at the surface of Ge1−xSnx shell (a) region A-54 at. % Sn; region B-9 at. % Sn and (b) region A-53 at. % Sn; region B-8 at. % Sn. All scale bars are 20 nm.
wi are the intensity and width of the ith fitted peak with emission centered at E0(i):29 4
I = I0 +
∑ i=1
Ai wi
π 2
2⎤ ⎡ ⎛ E − E0(i) ⎞ ⎥ ⎢ ⎟⎟ exp −2⎜⎜ ⎢ ⎝ wi ⎠ ⎥⎦ ⎣
Si(111) substrates were chosen to minimize substrate PL. To confirm that the PL signal was indeed due to the nanowires, PL spectra were recorded from areas of the sample without nanowires where adventitious epitaxial thin film deposition was observed; PL intensity from these areas without nanowires but E
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Figure 7. Low-magnification image of nanowire deposited with PSnCl4/PGeH4 = 2 × 10−5 showing Sn-enriched coherent defect region. (a) Lowmagnification TEM of coherent planar defect in Ge1−xSnx. (b) HRTEM showing continuous lattice fringes across defect with two regions of interest for EDS analysis, (c) Ge−K peak and (d) Sn−K peak of EDS spectra from regions 1 and 2 from (b). Region 1 shows 19 at. % Sn; region 2 shows 9 at. % Sn, which may be enhanced due to proximity of the probe to the defect. Scale bar in (a) is 200 nm; scale bar in (b) is 10 nm.
0, Figure 8c) which exhibits only a weak direct gap peak at ∼0.77 eV (1590 nm). This is consistent with results from previous studies showing that the indirect gap (L-valley) transition is quenched by surface states in as-grown Ge nanowires and the expectation that a 980 nm pump laser should produce minimal heating-induced red shifting.30 Focusing on the highest pump power spectrum for a PSnCl4/ PGeH4 = 2 × 10−5 sample (shown in Figure 8a,b), we see that the Ge direct gap (Γ-valley) peak wavelength occurs at 1784 nm (0.695 eV) compared to 1550 nm (0.8 eV) for bulk Ge and that the Ge indirect gap (L-valley) peak wavelength occurs at 1965 nm (0.631 eV) compared to 1880 nm (0.66 eV) for bulk Ge. This is consistent with the tendency of tensile strain to cause the Γ-valley transition energy to decrease faster than that of the L-valley in germanium.31−33 In particular, the measured axial strain in the PSnCl4/PGeH4 = 2 × 10−5 core-shell nanowires (4 at. % as measured by XRD) is 0.55%. Germanium with tensile strain of ∼0.5% exhibits an ∼0.05 eV and ∼0.1 eV
decrease in indirect-gap and direct-gap transition energies, respectively.32 This is quite consistent with the observed shifts of the Ge indirect- and direct-gap emission peaks in these nanowires. Furthermore, for a Sn content in Ge1−xSnx of x = 0.04, an ∼0.2 eV decrease in the direct-gap transition energy is expected.34 Consistent with this expectation of a decreased direct-gap transition energy, we observe a transition at 0.578 eV, which we attribute to the Ge1−xSnx direct-gap (cf. to 0.8 eV for pure Ge). The 0.550 eV peak extracted from the PL data and attributed to the Ge1−xSnx indirect-gap transition is also consistent with the fact that adding Sn to Ge causes the directgap transition energy to decrease more than it does the indirect-gap transition energy.35 On the basis of the extracted energies of the PL features, we ascribe to the Ge1−xSnx shell a Sn composition of 3−4 at. %,34 which is consistent with the analysis of XRD data summarized in Table 1. Photoluminescence data from PSnCl4/PGeH4 = 2 × 10−5 Ge1−xSnx nanowires demonstrated that emission wavelengths vary minimally with pump power greater than or equal to 27.8 F
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Figure 8. (a) Photoluminescence peak fitting and deconvolution for PSnCl4/PGeH4 = 2 × 10−5 core-shell nanowires at 60.0 kW cm−2 pump fluence; (b) photoluminescence of Ge/Ge1−xSnx core-shell nanowires (PSnCl4/PGeH4 = 2 × 10−5) at various laser pump powers; (c) pure Ge nanowire control sample at 27.8 kW cm−2 pump fluence; and (d) comparison of Ge nanowire reference sample (black trace from c) and PSnCl4/PGeH4 = 2 × 10−5 (blue trace from b) at 27.8 kW cm−2 pump fluence.
kW cm−2. We hypothesize that for powers less than 27.8 kW cm−2, the Ge1−xSnx emission is reduced due to nonradiative recombination because of the presence of defects on the shell surface, leaving dominant emission only from the strained Ge core. With greater pump power (Figure 8b), photogeneration of carriers and spillover from the core to the shell gives strong photon emission from the Ge1−xSnx shell. Figure 8d compares the PL spectrum measured from the core−shell and single component Ge nanowires at 27.8 kW cm−2 pump fluence. The Ge/Ge1−xSnx core-shell nanowire PL signal is orders of magnitude higher than that of pure (unstrained) Ge nanowires. In HRTEM, we observe local regions of coherent defects with significantly higher Sn composition (Figure 7). The coherent interfaces around the Sn-rich planar defects are not expected to contribute trap states that promote nonradiative recombination. We do not detect a PL feature attributable to the Sn-rich regions themselves, presumably because their very high Sn contents shift their emission wavelength to a value outside the range of our measurements (1400 nm < λ < 2400 nm).36
The PL data are closely linked to the Sn composition and the strain state of the core-shell Ge/Ge1−xSnx nanowires. Their PL emission wavelength is correlated to the core and shell lattice parameters as measured by XRD. The discrepancy between the XRD-measured lattice parameters and the EDS-measured compositions can be explained by the presence of Sn-rich precipitates and coherent defects. In summary, epitaxial Ge/Ge1−xSnx core-shell nanowires were synthesized by vapor−liquid−solid (VLS) growth. Indirect and direct bandgap transitions in Ge and Ge1−xSnx observed in these nanowires suggest a pure Ge core nanowire acting as a compliant substrate for subsequent growth of a Ge1−xSnx shell. The PL signal is strong even at room temperature and is orders of magnitude higher in Ge/Ge1−xSnx core-shell nanowires compared to control Ge nanowires, indicating that radiative recombination of photogenerated carriers in the Ge core is strongly enhanced in the core-shell wires compared to the situation for native oxide coated Ge nanowires. At the same time, significantly enhanced PL intensity is observed in the largely unstrained shells compared G
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to the tensile-strained Ge cores, indicating the promise of Ge1−xSnx nanowires for photonic applications. Interestingly, the core Ge acts as a compliant substrate for shell growth, adopting a lattice parameter in the vertical direction close to that of the Ge1−xSnx shell. This is similar to using a compliant substrate for two-dimensional Ge1−xSnx epitaxy. As a result, misfit dislocations are not observed in these nanowires, despite a large lattice mismatch. Placing the core Ge nanowire under tension enhances its PL by decreasing the energy of the Γ-valley relative to the L-valley. Furthermore, because both the direct and indirect gap transitions are strong in the Ge core, we conclude that carrier spillover from the core to the Ge1−xSnx shell will contribute to the observed strong room-temperature photoluminescence as long as the optical generation rate in the Ge core exceeds the rate of nonradiative recombination due to defects present on the shell outer surface.37 Experimental Section. Core−shell Ge/Ge1−xSnx nanowires were grown via VLS growth by CVD, using randomly dispersed 40 nm diameter colloidal Au nanoparticles on Si(111) and Ge(111) substrates in a three-step process (Figure 1). The first two steps were conducted under GeH4/H2 flow with PGeH4/PH2 = 0.0145 and included a 2 min nucleation step conducted at 375 °C followed by a 30 min growth segment at 300 °C for the Ge core nanowire. This produced ∼4 μm long nanowires with very uniform length.38,39 The 40 nm diameter Au colloids produce Ge core nanowires of ∼50 nm diameter. The last step in the CVD process was a 30 min growth segment conducted under GeH4/H2/SnCl4 flow at 300 °C with PGeH4/ PH2 = 0.0145 and PSnCl4/PGeH4 ranging from 0 to 3 × 10−5 with SnCl4 flow controlled by a calibrated needle valve. The SnCl4 precursor was introduced using H2 carrier gas through an electro-polished stainless steel bubbler (Strem) that was cooled to 1 °C. Total pressure during growth was 30 Torr. Growth of control Ge nanowires was conducted with only the first two steps, as reported previously.15,16 The as-deposited core-shell Ge−Ge1−xSnx nanowires grown on Si(111) substrates were probed at room temperature using micro-PL with a 980 nm diode laser focused to an ∼8 μm diameter spot size with a Mitutoyo 20× M Plan NIR microscope objective (NA = 0.4) in a surface normal pump/ collection geometry. A Stanford Research Systems SR830 Lockin Amplifier connected to an extended InGaAs photodiode cooled to −20 °C was used with a mechanical chopper at 300 Hz for phase sensitive detection of PL. The detector response was calibrated with an Ocean Optics tungsten halogen lamp modeled as a blackbody at 2960 K. SEM images were collected using a FEI Helios NanoLab 600i DualBeam Focused Ion Beam/Scanning Electron Microscope with 2 kV acceleration voltage and 43 pA beam current. TEM images were collected using an FEI Tecnai G2 F20 TEM operating at 200 kV. Nanowire samples were scraped onto holey carbon TEM grids (Electron Microscopy Sciences) with a dry stainless steel razor blade. XRD patterns were obtained using a PANalytical X’Pert using monochromated Cu Kα1 radiation via a hybrid X-ray mirror and 2 crystal Ge(220) 4bounce monochromator. High-resolution scans were also performed with a 3-bounce Ge(220) analyzer crystal in the diffracted beam path.
Letter
ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.nanolett.6b03316. Low magnification TEM images, details of elasticity model calculations applied to XRD data, TEM-EDS calibration standards, individual TEM-EDS spectra, PL spectra fit parameters and statistics, scanning transmission electron microscope (STEM) image of coherent defects (PDF)
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Author Contributions
A.C.M. and J.P.M. built and optimized the precursor delivery apparatus of the CVD reactor. A.C.M. performed all growth experiments and SEM characterization. C.S.F. and A.C.M. collected PL spectra. M.R.B. and A.C.M. performed XRD. A.C.M. and A.F.M. performed TEM analysis. A.C.M. and P.C.M. designed the experiments, and all authors contributed to writing the manuscript. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work is supported by National Science Foundation Division of Materials Research programs DMR-1206511 and DMR-1608977. Part of this work was performed at the Stanford Nano Shared Facilities. A.C.M. acknowledges financial support from a National Science Foundation Graduate Fellowship. C.S.F. acknowledges financial support from a Natural Sciences and Engineering Research Council of Canada fellowship.
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