Letter pubs.acs.org/NanoLett
Semiconductor-Oxide Heterostructured Nanowires Using Postgrowth Oxidation Jesper Wallentin,†,§ Martin Ek,‡ Neimantas Vainorious,† Kilian Mergenthaler,† Lars Samuelson,† Mats-Erik Pistol,† L. Reine Wallenberg,‡ and Magnus T. Borgström*,† †
Solid State Physics, Lund University, Box 118, S-221 00, Lund, Sweden Polymer and Materials Chemistry/nCHREM, Lund University, Box 124, S-221 00, Lund, Sweden
‡
S Supporting Information *
ABSTRACT: Semiconductor-oxide heterointerfaces have several electron volts high-charge carrier potential barriers, which may enable devices utilizing quantum confinement at room temperature. While a single heterointerface is easily formed by oxide deposition on a crystalline semiconductor, as in MOS transistors, the amorphous structure of most oxides inhibits epitaxy of a second semiconductor layer. Here, we overcome this limitation by separating epitaxy from oxidation, using postgrowth oxidation of AlP segments to create axial and core−shell semiconductor-oxide heterostructured nanowires. Complete epitaxial AlPInP nanowire structures were first grown in an oxygen-free environment. Subsequent exposure to air converted the AlP segments into amorphous aluminum oxide segments, leaving isolated InP segments in an oxide matrix. InP quantum dots formed on the nanowire sidewalls exhibit room temperature photoluminescence with small line widths (down to 15 meV) and high intensity. This optical performance, together with the control of heterostructure segment length, diameter, and position, opens up for optoelectrical applications at room temperature. KEYWORDS: Nanowire, MOVPE, III-V, quantum dots, oxide
D
or the overall morphology, creating monolithic structures of crystalline semiconductor segments separated by insulating, amorphous, oxide segments. PL measurements show promising RT characteristics with significantly smaller line widths (down to 15 meV) than recently reported colloidal quantum dots14,15 and similar intensity. NWs were epitaxially grown in an oxygen-free atmosphere using the vapor−liquid−solid technique with 40 nm gold catalysts (see Methods for details). First, axial InP−AlP−InP− AlP heterostructures were created by switching precursors during growth. After the entire structures were grown, the samples were taken out and exposed to air (Figure 1a). The samples were examined after growth and oxidation using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Postgrowth oxidation can be highly destructive in thin films,8 presumably due to volume expansion, but the open geometry of the NWs allowed them to remain morphologically intact (Figure 1b−d). The composition of the oxidized AlP, hereafter termed AlOx, could not be exactly determined, but measurements indicate a mixture of aluminum oxide and aluminum hydroxide (Supporting Information Section 1). The AlOx was
evices utilizing quantum confinement at room temperature (RT) require high charge carrier potential barriers. Self-assembled and NW-based quantum dots (QDs) show inhibited or reduced photoluminescence (PL) at RT,1−4 since thermally excited carriers can overcome the confining barrier, while colloidal QDs have shown complex temperature dependence.5 The heterointerface between semiconductors and oxides can be several electron volts high,6 but the amorphous nature of most oxides has hindered bottom-up synthesis of semiconductor-oxide heterostructures with quantum confinement in three dimensions. Top down processing with lithography and wet oxidation of lattice matched AlAs layers can be used to create oxide layers in relatively large devices.7 We have demonstrated the synthesis of AlP nanowires (NWs),8 where GaP shells prevented the spontaneous oxidation of AlP. In this Letter, we instead take advantage of postgrowth oxidation of AlP to create semiconductor-oxide heterostructured NWs. These NW heterostructures are highly lattice mismatched,9 which in thin films would lead to formation of propagating defects. However, it has already been shown that NWs can combine high quality III-V materials in an almost unlimited manner,10 as well as growth on silicon11 and graphene12 substrates, and recently hybrids of Si and IIIVs.13 Importantly here, the open geometry of the NWs allows oxidation without destruction of the semiconductor segments © 2013 American Chemical Society
Received: August 20, 2013 Published: November 6, 2013 5961
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Figure 1. Semiconductor-oxide heterostructured NWs by postgrowth oxidation of InP−AlP−InP−AlP NWs. (a) Drawing explaining the structure and the process steps. (b) SEM image of NWs with a long (6 min growth time, about 1 μm long) second InP segment (InP 2), (c) SEM image of NWs with a short (30 s growth time, about 50 nm long) second InP segment. (d) SEM image of InP−AlOx NW with multiple segments. The InP segments appear slightly thicker and brighter than the AlOx segments. All SEM images have a tilt 30° off normal.
Graded junctions have previously been observed in InAs-GaAs junctions due to In storage in the Au seed particle,19−21 and the asymmetry has been attributed to different dynamics in how the group III elements lower each other’s equilibrium concentrations.21 The AlxIn1−xP cones were less expected than the junction asymmetry. Somewhat similar cones were observed and attributed to strain, in InGaN-GaN NW junctions.22 No radial concentration gradients in the cones were observed (Figure 2g) and the boundary between the AlxIn1−xP and the AlOx shell was sharp. This indicates that each epitaxial layer was grown with a single-composition AlxIn1−xP, surrounded by an AlP shell. We note that the Al−In phase diagram shows a strong tendency for segregation23 and speculate that a combination of strain and liquid-phase segregation could explain our observations. However, we cannot rule out that the AlxIn1−xP with low In concentration oxidizes more readily, leaving a gradually thinner core of crystalline semiconductor. Further studies are underway to clarify this issue. Abrupt junctions are desirable to fully take advantage of the high energy barrier of the oxide. We therefore used growth stops between InP and AlP growth with only PH3 turned on in order to purge the seed particle of In before introducing Al.21 While this reduced the cone length to about 50 nm (Figure 2d,e), longer stops than one minute did not further reduce the cone size. Evidently, some In remained in the seed particle even after the growth stop and was released by the introduction of Al. The AlxIn1−xP cone length could possibly be further reduced by pulsing Al.21 Thus, we were not able to achieve isolated AlOx−InP−AlOx segments with sharp junctions in the axial direction.
determined to be amorphous due to a complete lack of crystalline reflections in the selected area electron diffraction patterns. For long segment lengths (about 1 μm) the InP→AlOx junctions were straight, while about 90% of the reverse AlOx→ InP junctions were kinked (Figure 1b). Such asymmetry is commonly observed during group III switching of heterostructured III−V NWs.16 However, when growing short (50 nm) InP or AlP segments we observed close to 100% straight AlOx→InP heterojunctions (Figure 1c), allowing growth of straight NWs with several AlOx and InP segments (Figure 1d). The heterojunctions were investigated using TEM (Figure 2). We mapped the spacings of the (111) planes perpendicular to the growth direction from high resolution TEM (HRTEM) images using geometrical phase analysis (GPA) (Figure 2g),17 and converted to the composition using Vegard’s law. Prior measurements of sharp NW heterostructures have shown that strain is efficiently relaxed in this direction,18 meaning that the spacing should accurately reflect the composition. The AlOx→InP junctions were slightly uneven (on the scale of about 5 monolayers) but otherwise abrupt, with measured (111) plane distances corresponding to the value of unstrained InP immediately after the onset of crystalline material. We did not find any indications that the AlOx was straining the semiconductor segments. The reverse InP→AlOx junctions, although straight for all segment lengths, were more complex. We observed axiocentric cones of ternary AlxIn1−xP (Figure 2bc), with gradually decreasing In concentrations and lengths which were about three times the NW diameters (150 nm). The crystal structure of the InP and the AlxIn1−xP was a mix of wurtzite and zincblende with a high density of stacking faults. 5962
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Figure 2. TEM of oxidized InP-AlP nanowire heterojunctions. (a) HRTEM of an amorphous AlOx to crystalline InP junction, (b) HAADF and (c) HRTEM images (stitched) of InP to AlOx junctions, displaying a cone of crystalline AlxIn1−xP. (d) HAADF and (e) HRTEM images of two InP to AlOx junctions, where a 1 min growth stop was used before AlP growth in order to reduce the In concentration in the seed particle. This reduced the AlxIn1−xP cone length by about 70%. (f) HRTEM, (g) (111) spacing map, and (h) HAADF of short (50 nm, 60 s growth time) InP segments preceded and followed by AlOx segments (as in Figures 1c and 3a,c). The growth direction is upward in all images.
Small amounts of radial (core−shell) growth often occurs on the NW sidewalls,24 which can be uneven such that quantum dots (QDs) form.25,26 After growth of short axial InP segments, TEM revealed hundreds of nanometer-sized QDs around the perimeter of the first AlOx segment (Figure 3b,c). No such dots were found on the second AlOx segment (Supporting Information Section 2). AlP NWs grown under these conditions have a high density of crystal stacking faults,8 which can act as preferential nucleation sites for radial growth.25,27 This vapor−solid growth is not susceptible to seed particle storage effects, which caused the graded axial junctions. Electron diffraction on some of the largest segments showed that these were indeed InP with a crystalline orientation almost parallel to that of the axial InP segment (Supporting Information Section 3). Thus, we were able to
synthesize isolated InP QDs in AlOx in the radial direction (Figure 3a). Compared with established QD materials systems, a key opportunity of our NWs is the high energy barrier of AlOx. In Figure 3d, we show the band diagram of an AlOx−InP−AlOx heterostructure, assuming that the AlOx is Al2O3 with a bandgap of 6.5 eV,6 and of two common QD systems. To study their optical properties we transferred NWs to gold-covered substrates and performed single NW PL at RT and low temperature (LT, T = 4.2 K), using continuous excitation with a slightly focused green laser (2.33 eV). For each NW, the LTPL exhibited several sharp peaks at different energies from 1.5 to 2 eV with the most and the brightest objects appearing in the middle of this range. Reference NWs grown without an InP segment did not exhibit such peaks. Assuming a highly simplified cubic infinite potential well model with the LT 5963
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Figure 3. TEM and PL measurements of radially grown InP segments in AlOx. (a) Schematic showing how small InP QDs grow radially on the first AlP segment, during growth of an axial InP segment, and become isolated in the AlOx upon oxidation. (b,c) STEM (HAADF) of InP QDs, seen as bright dots, in AlOx. No such dots were found in the AlOx above the axial InP segment (Supporting Information Section 3). (d) Figure comparing the band diagram of the InP/AlOx system with two common QD systems, as well as the energy (2.33 eV) of the laser used in our PL measurements. (e,f) Time trace of PL spectra of two different NWs, measured at low (T = 4.2 K) and room temperature, respectively. The integration times were 100 and 500 ms, respectively, and the excitation power density 2 W/cm2. (g,h) Average and single-shot spectra for the plots in (e,f), respectively.
As discussed above, the high energy barrier of the oxide holds promise of excellent RT optical performance. The highest brightness observed at LT and RT (different NWs) showed similar spectrally integrated intensities, about 1 × 105 counts/s (Supporting Information Section 5). With an estimated collection efficiency of 2.5% (Supporting Information Section 6), this corresponds to a QD intensity of 4 × 106 counts/s. These values are about 1 order of magnitude higher than NW QDs3 and also slightly higher than recent colloidal QDs,14,15 although precise comparisons are difficult due to different excitation and collection conditions. The single-shot and averaged spectra show RT line widths of 21 and 42 meV (Figure 3g,h), respectively, while the lowest observed value is 15 meV (Supporting Information Figure S4a). This is similar to NW-based3 and self-assembled QDs32 and significantly lower than the about 45 meV reported for colloidal QDs.33 Thus, while the QDs in our nonoptimized structures show much variation at this point, these results demonstrate the method’s potential. Another key opportunity of these structures is the ability to control the spatial positions, which is challenging with colloidal and self-assembled QDs. The diameters and lateral positions of the axial InP segments are governed by the Au seed particle (Figure 4a,b).34 Since the length and vertical position of the axial segments are determined by the switching during epitaxy, the size and position of the axial segments can be independently controlled in three dimensions. Controlling the radially grown QDs is more challenging. The size is related to the growth time and the radial growth rate. The QD density
bandgap (1.42 eV) and effective masses of InP the energy range corresponds to InP QDs in the size range 3 to 10 nm, consistent with the TEM observations. Thus, we assign the different peaks to highly quantized individual InP QDs with slightly varying size. In some NWs, we also observed signal around the InP bandgap, which could be due to the axially grown, 50 nm long, InP segment. The sharp, bright peaks allowed for integration times down to 20 ms, and we measured repeated spectra for the same NWs. A trace of PL spectra over time (Figure 3e) shows how the signal jumps between different positions, so-called spectral diffusion (SD),28 and dark periods, so-called blinking or intermittency.29 Blinking and SD are commonly observed in QDs 14,28−30 and are usually associated with trapped charges.14,28,31 The single shot spectrum in Figure 3g exhibits a LT line width of 1.3 meV, while SD broadened the integrated spectrum to 5.5 meV. With increasing excitation power, more lines were observed, the blinking and SD were faster, and both the peak intensities and the line widths increased (Supporting Information Section 4). Some QDs were dark only a few percent of the observation time at LT, although others were mostly dark. While we observed dozens of peaks in many NWs, the number of objects observed in PL was substantially less than in TEM (bright fraction on the order of 1 to 10%). One possible explanation is that the excitation due to the lack of a crystalline wetting layer is done resonantly and excites only those QDs that have a transition matching the laser energy. It is also possible that the uncontrolled oxidation led to nonradiative states forming at the semiconductor−oxide interface. 5964
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PH3, χPH3 = 6.3 × 10−3. Then, an InP segment was grown with lower TMI, χTMI = 3.2 × 10−6, and PH3, χPH3 = 3.1 × 10−3. Finally, a second AlP segment was grown the same way as the first one. The NWs investigated in Figures 2 and 3 were grown with indirect-bandgap GaP substrates and stubs to minimize disturbing optical signals. For the GaP−AlAs−InAs−AlAs heterostructured nanowires in Figure 4c, the PH3 was replaced with arsine (AsH3) for the top three segments. The molar fraction was χAsH3 = 2.2 × 10−4 for the AlAs segments and χAsH3 = 1.1 × 10−4 for the InAs segment. The transmission electron microscopy (TEM) analyses were performed on a JEOL 3000F operated at 300 kV. TEM samples were prepared by breaking off NWs from their substrates and transferring them to lacey carbon Cu grids by gently pressing the grids to the substrates. Local lattice parameter measurements by geometric phase analysis (GPA) were chosen for compositional analysis. The alternative method of energy dispersive X-ray spectroscopy would systematically have indicated too high Al contents, since X-rays from the surrounding AlOx shell would also have been detected. In addition, the isolated axial InP dots proved beam sensitive and gradually lost P on exposure, leaving behind In droplets. This necessitated an analysis method which required relatively low e− doses.
Figure 4. (a) TEM of axially grown InP segment in AlOx using 15 nm Au seed particles. Reducing the nanowire diameter by using smaller seed particles also reduced the AlxIn1−xP cone length. (b) SEM of position-controlled GaP−AlOx−InP−AlOx NWs grown from EBLdefined Au seed particles. Tilt 30°. (c) SEM of GaP−AlOx−InAs− AlOx NW, where the two AlOx segments were grown as AlAs instead of AlP. Tilt 30°.
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ASSOCIATED CONTENT
S Supporting Information *
Composition of AlOx. Distribution of QDs. Orientation of InP fragments and QDs. PL power dependence. Room temperature PL. Collection efficiency. This material is available free of charge via the Internet at http://pubs.acs.org.
and the vertical positions may be related to the crystal structure and stacking faults,25,27 as discussed above, whose position could be controlled by tuning the growth conditions.35 While we focused our studies on AlP, initial experiments indicate that postgrowth oxidation also works with the more common material AlAs (Figure 4c). The method presented here should be applicable to any oxygen-free NW growth technique, including molecular beam epitaxy (MBE), chemical beam epitaxy (CBE), and laser ablation techniques. The quality of the oxide could be improved by oxidizing with a controlled O2 partial pressure, which is not currently possible in our system. Optimizing the quality of the oxide will likely improve the optical properties. The method presented here expands the possible materials in NW heterostructures to include insulating, amorphous oxides with high energy barriers, and paves the way for optoelectronic devices operating at room temperature. Methods. Samples were prepared for NW growth by depositing 40 nm Au particles with an aerosol technique36 at a density of 0.5 μm−2 on InP or GaP (111)B substrates. The NWs were grown in a low-pressure (100 mbar) metal organic vapor phase epitaxy (MOVPE) system with a total flow of 6 l/ min using hydrogen as carrier gas. To desorb surface oxides, the samples were first annealed at 550 °C (InP) or 650 °C (GaP) for 10 min under a PH3/H2 gas mixture. Then, the temperature was lowered to 460 °C. Growth started with a NW stub of the same material as the substrate, which was grown for 1 min. Trimethylindium (TMIn) with molar fraction χTMI = 9.5 × 10−6, and phosphine (PH3) with χPH3 = 6.3 × 10−3, were used as precursors for InP with HCl (χHCl = 1.3 × 10−5) in situ etching to prevent radial growth.37 Trimethylgallium (TMGa), χTMG = 1.8 × 10−5, and PH3, χPH3 = 6.3 × 10−3, were used as precursors for GaP. Next, the first AlP segment was grown for 3 min using trimethylaluminum (TMAl), χTMA = 2.8 × 10−5, and
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Present Address §
Institute for X-ray Physics, University of Göttingen, FriedrichHund-Platz 1, 37077 Göttingen, Germany. Author Contributions
J.W. and M.T.B. grew the samples and wrote the manuscript. M.E. and L.R.W. performed TEM analysis. J.W., N.V., K.M., and M.E.P. performed optical measurements. M.E.P., L.R.W., L.S., and M.T.B. supervised the project. All authors discussed the results and commented the manuscript. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was performed within the Strategic Focus Area Nanometer Structure Consortium at Lund University (nmC@ LU) and was supported by the Swedish Research Council (Vetenskapsrådet), by the Knut and Alice Wallenberg Foundation, and by the Swedish Energy Agency.
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