Letter pubs.acs.org/NanoLett
Epitaxial Growth of GaN Nanowires with High Structural Perfection on a Metallic TiN Film M. Wölz, C. Hauswald, T. Flissikowski, T. Gotschke, S. Fernández-Garrido, O. Brandt, H. T. Grahn, L. Geelhaar,* and H. Riechert Paul-Drude-Institut für Festkörperelektronik, Hausvogteiplatz 5−7, 10117 Berlin, Germany ABSTRACT: Vertical GaN nanowires are grown in a self-induced way on a sputtered Ti film by plasma-assisted molecular beam epitaxy. Both in situ electron diffraction and ex situ ellipsometry show that Ti is converted to TiN upon exposure of the surface to the N plasma. In addition, the ellipsometric data demonstrate this TiN film to be metallic. The diffraction data evidence that the GaN nanowires have a strict epitaxial relationship to this film. Photoluminescence spectroscopy of the GaN nanowires shows excitonic transitions virtually identical in spectral position, line width, and decay time to those of state-of-the-art GaN nanowires grown on Si. Therefore, the crystalline quality of the GaN nanowires grown on metallic TiN and on Si is equivalent. The freedom to employ metallic substrates for the epitaxial growth of semiconductor nanowires in high structural quality may enable novel applications that benefit from the associated high thermal and electrical conductivity as well as optical reflectivity. KEYWORDS: Semiconductor, metal, heteroepitaxy, self-induced growth, nanorod, nanocolumn
T
order to exploit the benefits of metallic substrates for LEDs, that is, the excellent electrical and thermal conductivity as well as optical reflectivity, ensembles of vertical (In,Ga)N/GaN NWs with LED structure have also been transferred after growth to Cu substrates.17 Nevertheless, for most applications direct growth is preferred over transfer processing. Moreover, in none of the previous examples mentioned in this paragraph the growth of the semiconductor NWs was epitaxial. In this Letter, we demonstrate the direct epitaxial growth of GaN nanowires (NWs) with high crystal quality on a metallic TiN film. Our approach is based on the self-induced growth of GaN NWs by molecular beam epitaxy (MBE), by which the NWs form spontaneously18,19 and largely independent of the substrate [see ref 20 and references therein]. Besides structural mismatch, another hurdle to be overcome for epitaxial growth are interfacial reactions that are likely to occur for chemically dissimilar materials, in particular semiconductors on metals.21,22 In this framework, it is advantageous that GaN is chemically fairly inert. Our choice of substrate has been guided by the fact that TiN is a refractory material that is used as technical ceramics and is stable not only against GaN but also its constituents even at temperatures exceeding those required for GaN NW growth.23 At the same time, TiN exhibits metallic conductivity and forms an ohmic contact to GaN.24 As a matter of fact, the growth of GaN NWs on TiN by MBE has been noted before when using TiN as a mask for selective-area growth but was considered parasitic and not investigated
he bottom-up growth of free-standing nanowires (NWs) has opened up new opportunities for the epitaxial combination of dissimilar materials because strain induced by mismatch in lattice constants or thermal expansion coefficients can elastically relax at the free NW sidewalls.1,2 Moreover, even if dislocation lines form they are likely to be constrained at the interface between NW and substrate3 or to terminate at the NW sidewalls,4,5 thus rendering the upper part of the NWs defect-free. Early on, these advantages of NW epitaxy have been employed for the monolithic integration of III−V semiconductors on Si substrates,6 leading to the demonstration of several devices including light-emitting diodes (LEDs), lasers, photodiodes, and transistors.7−12 More recently, the epitaxial growth of semiconductor NWs on graphene has been pursued, motivated by the high conductivity of this kind of substrate, its mechanical flexibility, and its potentially low cost compared to single-crystalline semiconductors [see ref 13 and references therein]. Even less explored is the direct growth of semiconductor NWs on substrates made of conventional metals. Si NWs were grown on stainless steel foil and processed into solar cells.14 This approach is highly attractive because it could pave the way to roll-to-roll fabrication. Hence, the underlying growth mechanisms have been explored for several material combinations.15 However, in all these cases the nanowires grew as a grass-like ensemble without any preferential orientation, which thwarts both the controlled growth of heterostructures along the NW axis and subsequent device processing. Regular arrays of CdS NWs were grown on Al foil by converting part of this foil into a highly periodic anodic alumina membrane serving as a template for NW synthesis.16 Still, this technique is limited to this particular type of metal foil and template-guided growth. In © XXXX American Chemical Society
Received: January 21, 2015 Revised: May 7, 2015
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DOI: 10.1021/acs.nanolett.5b00251 Nano Lett. XXXX, XXX, XXX−XXX
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Figure 1. Evolution of the RHEED pattern during nitridation of the Ti film and GaN deposition. The shadow edge is at the top margin of the images. (a) The pattern of the Ti film prior to exposure to Ga or N indicates a smooth and well-oriented film. (b) The pattern of the Ti film shortly after exposure to Ga and N at 850 °C suggests the formation of δ-TiN. (c) The pattern after 90 min of exposure of the TiN film to Ga and N at 750 °C evidences the formation of epitaxial GaN.
substrate temperature to 850 °C as measured by a pyrometer and opened the Ga and N shutters simultaneously. (The pyrometer reading was calibrated for Si with an emissivity of 0.7. The emissivity of Ti is about 0.6 and the pyrometer reading hence underestimates the Ti surface temperature.) The streaky RHEED pattern of the Ti film presented in Figure 1a demonstrated that the sputtered film was smooth and singlecrystalline with an in-plane orientation such that αTi⟨112̅0⟩∥α-Al2O3⟨11̅00⟩. Upon opening the shutters, the pattern changed and rapidly evolved into the transmission pattern shown in Figure 1b. This pattern indicates the formation of a cubic crystal with a [111] out-of-plane orientation and the ⟨11̅0⟩ in-plane orientation parallel to the ones mentioned above. Diffraction spots arising from rotational twins and multiple diffraction are also present. The cubic symmetry and the lattice constant, which can be deduced from the pattern, are consistent with δ-TiN formed due to the nitridation of the Ti film by the N plasma. The change from a streaky to a spotty RHEED pattern suggests a simultaneous roughening of the surface. The characterization of reference samples by atomic force microscopy revealed that the Ti film consists of irregularly shaped smooth plateaus separated by trenches. The smooth plateaus are consistent with the streaky RHEED pattern, but because of the trenches the root-meansquare roughness of the Ti film is 5 nm. Exposure to active N results in a randomly corrugated surface without any plateaus, again consistent with the RHEED data, and at the same time the root-mean-square roughness increases only marginally to 6 nm. Coming back to the description of the main experiment, at the chosen high temperature the desorption of Ga is high and nucleation of GaN cannot occur. We approached the nucleation conditions for GaN NWs on the TiN film by gradually reducing the substrate temperature by −1 K/min under continuous Ga and N flux. The diffraction spots in RHEED, which reflect the hexagonal symmetry characteristic for α-GaN, were observed when the indicated temperature reached 750 °C. The temperature was then kept constant, and the final RHEED pattern after 90 min of growth is shown in Figure 1c. The diffraction spots evidence the formation of epitaxial GaN with an orientation-relationship such that αGaN(0001)⟨112̅0⟩∥δ-TiN(111)⟨11̅0⟩. The uniform crystallographic orientation can be deduced from the narrow rotational distribution of the spots around the specular beam and the absence of spots from other azimuths. The scanning electron micrographs in Figure 2a,b show that the RHEED pattern in Figure 1c originated from GaN NWs
further.25 To some extent related is also a report in which the growth of GaN NWs by metal organic vapor phase epitaxy on sapphire and GaN was induced by a thin film of Ti.26 However, in that case it was not shown that the metal film really acted as a substrate. Here, we present a detailed study of the NW nucleation conditions and the resulting GaN crystal quality. The GaN NWs were grown by plasma-assisted MBE on the TiN film at a substrate temperature of 750 °C with a Ga flux corresponding to a planar GaN growth rate of 2.4 nm/min and a N flux corresponding to 13 nm/min. As a reference, we used a GaN NW ensemble grown self-induced on Si(111) at 820 °C with source fluxes of 4.7 Ga and 13 nm/min N for 6 h. The high substrate temperature was chosen to minimize coalescence and thus to maintain the high crystal quality of the GaN NWs. Under these conditions, the loss of Ga atoms from the surface by desorption is significant, and a higher Ga flux is required compared to the sample grown on the TiN film. Variable-angle spectroscopic ellipsometry was performed at room temperature in a Sopra GES5E ellipsometer. For PL measurements, both samples were mounted into a microscope cryostat. Continuous-wave μ-PL was carried out exciting the sample with the 325 nm line of a Kimmon He−Cd laser through a 15× microscope objective resulting in a diameter of 3 μm. The luminescence was collected by the same objective, spectrally dispersed by a 2400 l/mm grating providing a spectral resolution 0.25 meV and detected by either a charged coupled device or an (In,Ga)As array. Time-resolved PL measurements were performed by exciting the samples with the second harmonic of pulses obtained from an optical parametric oscillator (OPO) synchronously pumped by a femtosecond Ti/ sapphire laser, which itself was pumped by a frequency doubled Nd/YVO4 laser. The pulse length after the OPO was about 200 fs, and the repetition rate of the laser system was 76 MHz. The frequency-doubled laser pulses have a wavelength of 325 nm and were focused by a 10× objective onto the samples resulting in a minimal spot size of 3 × 11 μm2 and a fluence of 0.8 μJ/ cm2. The luminescence signal from the samples was spectrally dispersed by a monochromator providing a spectral resolution of 4 meV and detected by a streak camera with a temporal resolution of 50 ps. For sample growth, a Ti film of 1 μm thickness was sputtered onto a polished sapphire [α-Al2O3(0001)] wafer, and all further processes were carried out by plasma-assisted MBE under in situ monitoring by reflection high-energy electron diffraction (RHEED). The nucleation and growth of GaN NWs requires a high surface temperature and a N-to-Ga adatom density ratio (local V/III ratio) above unity.19,27 Thus, we first set the B
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Figure 2. Scanning electron micrographs of the GaN NW grown on TiN in bird’s eye view for (a) low and (b) high magnification.
that are vertically aligned on the substrate. The NWs are on average 200 nm long (the length of well developed NWs is 300 ± 50 nm), typically 20−40 nm in diameter, and their density is 8 × 108 cm−2. Together with the [0001] orientation evident from the RHEED pattern in Figure 1c, these micrographs confirm that the NWs grow self-induced in the c-direction of GaN, as universally observed in MBE regardless of the substrate.20 A comprehensive growth study that is beyond the scope of this letter and will be reported elsewhere revealed that in general nucleation and growth of these NWs follow the same behavior as found on Si.28 The fairly wide dispersion in length observed here is the result of different nucleation times for individual NWs. More homogeneous NW ensembles can be obtained by growing for longer times or at lower substrate temperature. Also, length, diameter, and density of the NWs can be adjusted by the growth parameters. The as-grown wafer has a golden appearance as expected for TiN. X-ray diffraction ω-2θ scans performed on the as-grown sample29 (not shown) revealed the presence of not only TiN(111) but also TiO(111) as a result of the reaction of Ti with Al2O3 at the GaN growth temperature.30 The RHEED pattern does not permit to distinguish TiN and TiO, which have an identical crystal structure and similar lattice constant. We expect, however, that TiO forms at the interface of Ti and Al2O3 as observed, for example, by Morgan et al.,31 and TiN at the surface due to the reaction of Ti with the N plasma. In fact, the Raman spectra (not shown) of the as-grown wafer exhibit the characteristic first- and second-order modes of TiN.32 A contact-less method to experimentally investigate the electronic properties of the sample surface is variable-angle spectroscopic ellipsometry. Figure 3a,b shows the real (⟨ε1⟩) and imaginary (⟨ε2⟩) part of the pseudodielectric function of our sample, respectively. Data obtained at different angles were essentially identical, indicating that the contribution of the GaN NWs was negligible due to their low density. The behavior of ⟨ε1⟩ and ⟨ε2⟩ at low photon energies is characteristic for a metal. Moreover, our data are in good agreement with measurements on stoichiometric TiN from the literature.33,34 The exposure of the sputtered Ti film to the N plasma is thus an effective way to form metallic TiN suitable for the nucleation and growth of GaN NWs. As an indicator of the crystal quality of the GaN NWs grown on TiN, we employ low-temperature photoluminescence (PL) spectroscopy and compare the results to a state-of-the-art GaN NW reference sample grown on a crystalline Si substrate. The reference NWs had an average length of 2.7 μm, a mean diameter of 105 nm, and a density of 2 × 109cm−2. Figure 4 displays the normalized μ-PL spectra of the two samples
Figure 3. (a) Real and (b) imaginary part of the pseudodielectric function obtained by spectroscopic ellipsometry for the GaN NWs on TiN (blue dots) and the 65 nm thick TiN film (solid, green line) studied by Langereis et al.33 The dashed, orange line represents the pseudodielectric function of TiN obtained by electron energy loss spectroscopy by Pflüger et al.34
Figure 4. Normalized low-temperature (10 K) μ-PL spectra of the GaN NWs grown on TiN and Si under the same excitation conditions. Both samples are dominated by the recombination of A excitons bound to neutral donors [(D0,XA)], but also acceptor-bound [(A0, X)] and free exciton (XA) transitions are observed as shown in the inset.
obtained at 10 K and with an excitation density of 1 W/cm2. The spectra are dominated by the recombination of A excitons bound to neutral O and Si donors at 3.4713 [(O0,XA)] and 3.4721 eV [(Si0,XA)], respectively, that is, at essentially identical energies as for free-standing GaN.35,36 The recombination of free A excitons (XA) is visible at 3.478 eV. The inset of Figure 4 depicts the near-band-edge PL of the GaN NWs on TiN on a linear scale. Additional transitions from excitons bound to neutral acceptors [(A0,XA)] at 3.467 eV and B excitons bound to neutral donors [(D0, XB)] at 3.475 eV are evident on this expanded scale. The full width at half-maximum of the dominant donor-bound exciton transitions is about 1 meV and virtually identical for both samples. This value is among the lowest reported for such GaN NW ensembles20 and implies a high structural perfection. In the energy range of 3.40 to 3.43 eV, a set of sharp lines is observed, which is attributed to the radiative recombination of C
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recombination channel, since their mean NW diameter differs by a factor of 3. We have recently identified the coupling of bound exciton states as one mechanism giving rise to a biexponential decay of the (D0,XA) transition in GaN NWs.44 As detailed by Hauswald et al.,44 the short component of the biexponential decay approximately reflects the actual effective lifetime of the (D0,XA) complex. Within the frame of this interpretation, the lifetimes for the two samples are comparable. The lifetimes thus obtained are, however, shorter than the radiative lifetime of the (D0,XA) transition of at least 1 ns.45,46 Hence, for both samples a dominating nonradiative channel seems to be active. For samples grown on Si substrates, we have deduced that point defects are the most likely origin of the fast decay.47 Most importantly, here we did not find any indication for an additional nonradiative decay channel due to the growth on the metallic TiN. Our time-resolved PL measurements thus prove that the GaN NWs grown on TiN are of comparable crystal quality as GaN NWs grown on Si(111). In conclusion, we have demonstrated that GaN NWs of high crystal quality can be epitaxially grown on a metallic TiN film. This achievement in materials synthesis may pave the way for innovative applications that make use of the additional freedom in the choice of substrate properties, in particular with respect to electrical as well as thermal conductivity and optical reflectivity.
Figure 5. Time-resolved PL transients of the (D0, X) transition of the GaN NWs on TiN (triangles) and on Si (dots) at 6 K. The solid lines represent a biexponential fit to the data, and the lifetimes extracted from these fits are displayed next to the transients. The dashed line indicates the system response function of our setup.
excitons bound to I1 stacking faults in GaN.37 This type of stacking fault was shown to be induced by the coalescence of GaN NWs.38 However, the scanning electron micrographs shown in Figure 2 visualize that coalescence of adjacent NWs is unlikely to occur for the sample grown on TiN. We speculate that stacking faults may also form during the nucleation phase of GaN on TiN. Detailed microscopic investigations will be required to clarify the nature of the interface between the nitridated Ti film and the GaN NWs and to detect the presence of structural defects such as the stacking faults evident from Figure 4. In any case, the intensity of these PL lines is in both samples 2 orders of magnitude lower than that of the donorbound exciton transitions. Since PL is extremely sensitive to radiative defects like stacking faults, the stacking fault density must thus be very low. No PL signal was detected in a spectral range between 1.8 and 3.40 eV. Using a different setup, we have also acquired spectra in the infrared region including the energy at which emission from Ti2+ in GaN is expected (1.19 eV).39 We did not observe any emission related to Ti or other deep centers in these spectra. However, the potential incorporation of Ti, while not resulting in a detectable radiative transition, may lead to the formation of nonradiative centers in the GaN NWs. In addition, it was recently suggested that the bound exciton decay of thin GaN NWs is governed by nonradiative surface recombination even at low temperatures.40 The GaN NWs grown on TiN are only 20 to 40 nm thick and may thus be influenced by this process. To detect possible nonradiative decay channels, we performed time-resolved PL measurements. Figure 5 shows the PL transients of the (D0 ,XA) transition for both samples integrated spectrally over a 5 meV broad window centered at the maximum PL intensity. Because the decay is not monoexponential, we have used a biexponential decay to fit the data. Surface recombination in conjunction with the broad diameter distribution has been reported to result in a nonexponential PL decay of both ZnO and GaN NW ensembles.40−43 However, the decay times obtained in the present case (cf. Figure 5) are similar for the two samples under investigation, ruling out surface recombination as the dominant
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS We are grateful to C. Herrmann, H.-P. Schönherr, and C. Stemmler for the maintenance of the MBE system, A.-K. Bluhm for the scanning electron micrographs, M. Ramsteiner for his help with Raman measurements, D. van Treeck for the atomic force microscopy characterization, and R. Goldhahn for fruitful discussions of the ellipsometric measurements. Financial support of this work by the Leibniz-Gemeinschaft under Grant SAW-2013-PDI-2 is gratefully acknowledged.
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