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In this article, the authors have investigated the properties of the popular native defects in nitrogen-doped ZnO microrod samples grown by the chemical vapor transport method. Excellent crystalline quality has been confirmed in the samples. Optical signatures of zinc interstitials and zinc vacancies have been observed by employing Raman and variable temperature photoluminescence. By tuning the flow rate of oxidant (nitrous oxide) during growth, the concentration of zinc interstitials and vacancies can be modified. When the flow rate of the nitrous oxide is high, the zinc interstitials can be suppressed while the zinc vacancy-related shallow acceptors can be enhanced. These are both beneficial to the realization and further enhancement of p-type conductivity in ZnO material. This study provides a good understanding of the properties of the native point defects in nitrogen-doped ZnO microrods and also offers a simple way to control the defects towards p-type direction.
© 2016 Optical Society of America
1. Introduction
Over the past few years, wurtzite ZnO material has attracted huge interest due to its potential applications in short wavelength optoelectronic devices. Thanks to its high exciton binding energy, highly efficient light emission from excitonic recombination is expected. Furthermore, the possibility to fabricate a variety of nanostructures adds to the scientific and technological prospects of ZnO material [1]. In order to fabricate ZnO based optoelectronic devices and to enhance the device performance, reliable p-type doping is indispensable. Unfortunately, the reproducibility of p-type doping in ZnO is a major issue partially due to the self-compensation by native defects, such as zinc interstitial (Zni) and oxygen vacancy (VO). For instance, the Zni is generally regarded as a shallow donor with its energy level of 30 meV below the conduction band minimum (CBM) [2], and contributes to the n-type conductivity of ZnO [3]. Meanwhile, another intrinsic point defect zinc vacancy (VZn) has been suggested to be the dominant compensating acceptor in ZnO [4], although its energy position and optical fingerprints remain highly controversial [5–10].
Researchers have made great efforts to obtain stable p-type ZnO by using group I elements Li [11] and Na [12], and group V elements N [13], P [14], As [15], and Sb [16]. Among all these elements, N is widely considered the most promising candidate due to the similar ionic radius compared to O and availability of gas sources such as N2O, NO, NO2, and NH3. The possibility of forming p-type ZnO by substituting O site with N (NO) was first proposed in 1983 [17]. However, the origins of the shallow acceptor state in group V elements doped p-type ZnO are still being debated. The energy level of NO was determined to be 170-200 meV above the valence band maximum (VBM) in 2002 [13], but recent calculations have shown its metastable property [18] with an exceedingly high ionization energy of 1.3 eV [19], supported by the experimental results of Huang et al [20]. In their work, the red emission at above 700 nm has been proposed to relate to nitrogen deep acceptors created by post-growth annealing in ammonia. Similar results are applicable for other group V elements whose substituting defects (PO, AsO, and SbO) have very high acceptor-ionization energy [21]. In order to explain the experimentally measured p-type conductivity, people have proposed a new mechanism including two VZn and an antisite forming PZn−2VZn, AsZn−2VZn, and SbZn−2VZn, which have energy levels of around 150 meV above the VBM [21,22]. Lately, the shallow acceptor state of N-doped ZnO was proposed to originate from the defect complex VZn-NO, which actually evolved from the metastable double donor complex NZn-VO [23]. Furthermore, with positron annihilation spectroscopy, it has been reported that the introduction of N impurities into ZnO leads to the formation of stable vacancy clusters [24].
Nevertheless, we cannot neglect the roles of VZn and Zni on enhancing and compensating the p-type conductivity. At this stage, the precise control of VZn and Zni as well as nitrogen dopants opens up a new path to improve the p-type quality of ZnO. In a previous paper [25], we confirmed the formation of a shallow acceptor state in vertically aligned N-doped ZnO microrods (MRs), which were homoepitaxially grown on a high-quality ZnO template via chemical vapor transport (CVT) method, with N2O employed as both N and O precursors. The acceptor states have been assigned to complexes involving the intrinsic vacancy and nitrogen. While nitrogen here is an indispensable part to forming or stablizing the complex acceptor [23,24]. Meanwhile, optical signatures of VZn and Zni have been proposed. In this paper, with the purpose of intentionally controlling VZn and Zni, we will show and discuss the properties of a series of N-doped ZnO MR array samples grown at different flow rates of N2O. The change of the oxidant flow rate is the most convenient way to tune the chemical potentials, affecting the formation energy of native defects. According to the characterizations, shallow acceptors and compensating donors could be modified by changing the flow rate of N2O. We have achieved a suppression of donors and an enhancement of acceptors in samples grown with high flow rate of N2O.
2. Experiments
The growth of N-doped ZnO MR arrays was carried out via CVT method without employing any catalysts. The substrate was a 2-μm-thick high-quality ZnO template on sapphire grown by metal-organic chemical vapor deposition. N2 was used as the carrier gas, while N2O was employed as both N and O precursors. During the growth process, the pressure in the reactor was maintained at 3 kPa. The detail of the procedures and conditions has been reported elsewhere [25]. Four samples grown with different N2O flow rates of 1, 2, 3, and 4 standard cubic centimeter per minute (SCCM), were to be discussed in this paper and marked as samples A, B, C, and D, respectively.
High-resolution X-ray diffraction (XRD, Philips X'pert Pro diffractometer equipped with a Cu Kα x-ray source) was used to investigate the crystallinity of the samples. In addition, the vibrational properties of the four samples were recorded by a Raman system (JOBIN YVON HR800) in the backscattering geometry with 514 nm radiation at room temperature (RT). The polarized laser was perpendicular to the surface of samples with a spot size about 5 μm, and the power was 5 mW. The chemical configuration of the elements was determined by X-ray photoelectron spectrometry (XPS) with an Al Kα x-ray monochromatic source at 1486.6 eV. An X-band electron paramagnetic resonance (EPR) measurement was done to trace the defects by a Bruker EMX-10/12 spectrometer operating near 9.49 GHz at 110 K, using a rectangular TE102 resonator. The samples were cut into 5*2 mm2 and the magnetic field was perpendicular to the c axis of the ZnO MRs. For analyzing the optical properties, and especially the energy position of intrinsic zinc defects in the band gap, temperature-dependent (TD) PL spectra were recorded using a transverse magnetic polarized He–Cd laser at the wavelength of 325 nm with an angle of incidence about 15°. The laser spot size on the sample was 1.8 mm and the power was 40 mW.
3. Results and discussion
3.1. Crystallinity
The samples show similar morphologies as presented in Ref. 25. Different N2O flow rates do not change the morphologies, and a high-density N-doped ZnO MRs array with a hexagonal symmetry is vertically aligned on the substrate. The MRs have a uniform length of about a few micrometers and an average diameter of 1 μm. Figure 1(a) shows the XRD patterns of the four samples. The distinct ZnO (0002) diffraction peaks are observed and the weak peaks at 41.68° are reflections coming from the sapphire substrate. The shoulder of the main diffraction peak is believed to be from the ZnO template, which is strained on sapphire substrate due to lattice mismatch. Another significant observation from Fig. 1(a) is that the diffraction peaks of the ZnO MRs shift to higher angles (around 34.5°) compared with the (0002) peak of ZnO (34.42°) and the (222) peak of Zn3N2 (31.66°) [26], as shown in Table 1. Similar behavior was observed in N-doped ZnO thin films synthesized by chemical vapor deposition [27], indicating a reduction in the lattice constant and that no nitride phase or oxynitride alloy formation occurs in the ZnO MRs with N incorporation. One reason for the shift is that Zn-N bond lengths are somewhat shorter than Zn-O bond lengths [28]. However, the difference between the bond length of Zn-N and Zn-O is relatively small. It has been reported that the compressive stress randomly distributed in the lattice, which is caused by excessive N, can lead to an additional collapse of Zn tetrahedrons around NO [29]. Such a collapse of the lattice could be the main reason for the (0002) peak shift. In addition, the rocking curves around ZnO (0002) diffraction peaks are illustrated in Fig. 1(b), and the full width at half maximum (FWHM) values of samples A, B, C, and D are 0.1103°, 0.1040°, 0.0859°, and 0.0916°, respectively. All these results reveal that the N-doped ZnO MRs grow along the c axis with a well-ordered wurtzite structure.
Fig. 1 (a) XRD patterns of the MR samples. (b) The rocking curves around ZnO (0002) diffraction peak of N-doped ZnO MRs.
Table 1. The ZnO (0002) diffraction peak in the XRD patterns of our samples compared with the previous data in literature.
3.2. Vibrational properties
Raman analysis in back scattering geometry was also employed to study the structural properties of the samples, with incident light parallel to the c axis of the N-doped ZnO MRs. Wurtzite ZnO belongs to the space group C46v with two formula units in the primitive cell. The zone-center optical phonons can be classified according to the following irreducible representations: , where E2 modes are nonpolar and Raman active only. Figure 2(a) shows the Raman spectra recorded at room temperature with a resolution of 0.5 cm−1. The peaks located at 99, 332 and 438 cm−1 are classical ZnO modes of low frequency E2 (E2(low)), 2E2(M), and high frequency E2 (E2(high)), respectively. Furthermore, additional modes (AMs) located at around 276, 510, 582, and 644 cm−1 are all observed in the Raman spectra of the four samples, which have been found in various N-doped ZnO materials. They were previously assigned to local vibration modes due to NO [30], and it was reported that the intensity of the peaks at 272 cm−1 increases with N concentration [31]. However, these AMs were also observed in the Raman spectra of Fe-, Sb-, and Al-doped ZnO thin films [32]. And the AMs were consequently attributed to a disorder-induced relaxation of the translational crystal symmetry that allows the observation of the silent B modes [33]. Lately, the mode at 274 cm−1 was proposed to relate to the intrinsic point defect Zni or small Zni clusters based on precise logical inference [34,35], and it was reported that the formation energy of Zni in ZnO decreases with N incorporation by using ab initio density functional calculations [36]. As for the other three modes at 510, 582, and 644 cm−1, they usually emerge simultaneously with the mode at 276 cm−1 [30–33]. The intensity ratios of AMs over E2(low) as a function of the flow rate of N2O are illustrated in Fig. 2(b). It is found that the other three AMs follow the same trend as the mode 276 cm−1. Therefore, we attribute the AMs in our N-doped ZnO MRs to the vibrations related to intrinsic defect Zni, which is generally regarded as a shallow donor [3].
Fig. 2 (a) Raman spectra of the MR samples recorded at RT. (b) The intensity ratios of AMs and E2(low) as a function of the flow rates of N2O.
As shown in Fig. 2(b), the intensity ratio of the AMs over E2(low) firstly increases as the flow rate of N2O increases from 1 SCCM to 2 SCCM, then decreases and reaches the minimum value when the flow rate of N2O increases eventually to 4 SCCM. In order to explain why this occurs, we review the process of growth. The N2O was employed as both N and O precursors and the pressure in the reactor was maintained at the same value for growing different samples. When we increased the flow rate of N2O from 1 to 2 SCCM, the flow velocity of O atoms increased. Then the reacting time of O atoms in the reactor decreased. With the flow rates of N2O continuously increasing to 4 SCCM, more and more O atoms participate in the reaction and counteract the effect of reacting time. It is well known that the intrinsic donor defects Zni and VO in ZnO have higher formation energy during O-rich equilibrium growth [6,7]. Hence, from the Raman analysis, we suppose that the Zni related defects in the N-doped ZnO MRs could be suppressed at higher flow rates of N2O.
3.3. XPS analysis
In order to verify our hypothesis and investigate the chemical configuration of the elements, XPS was performed on samples A and D. Figure 3 shows the O 1s and Zn L3M45M45 auger lines with the binding energy scale calibrated by the C 1s peak at 285 eV. The typical asymmetric O 1s peaks in Fig. 3(a) can be consistently fitted by three nearly Gaussian components, centered at 530.2, 531.7 ± 0.1, and 532.4 ± 0.1 eV, respectively. Undoubtedly, the main component on the low binding energy side at 530.2 eV can be attributed to the O-Zn bond. And the component located with high binding energy at 532.4 ± 0.1 eV is usually attributed to chemisorbed or dissociated OH or O species on the surface of ZnO material, such as adsorbed H2O or O2 [37]. As for the component with the medium binding energy of 531.7 ± 0.1 eV, it is associated with O2- ions in the O deficient regions within the ZnO lattice, and the intensity of this component may be partially connected with the variations of the concentration of intrinsic defect VO [38,39]. For zinc interstitials, Awan et al [40] and Li et al [41] have reported that zinc ions at interstitial and lattice sites can be characterized by a different binding energy in the L3M45M45 auger lines (480 – 510 eV). Therefore, we have measured the LMM lines, which are shown in Fig. 3(b). The spectra have been fitted by two Lorentz-Gaussian components, centered at 495.3 ± 0.1 and 499.1 eV, which were attributed to Zni and lattice Zn, respectively [41]. Table 2 summarizes the percentage of the relative intensities of the three components of O 1s peaks and the two components of Zn LMM peaks. From sample A to sample D, the relative intensities of the components at 531.7 ± 0.1 and 495.3 ± 0.1 eV decrease, which suggests that the VO and Zni defects in the N-doped ZnO MRs grown at higher N2O flow rate are suppressed, confirming the Raman results.
Table 2. Relative intensities of the three components of O 1s peak and the two components of Zn LMM peak in sample A and D
3.4. EPR analysis
The change of intrinsic donors has been also reflected by EPR measurement. Figure 4 shows the EPR characterizations of samples A and D, respectively. Two obvious resonances at g ~1.955-1.957 and 2.002 could be detected. The g ~1.955-1.957 signal has been associated with an unpaired electron trapped on an oxygen vacancy [42] or on zinc interstitial related shallow donor [43]. Either attribution is valid since the strength of the signal weakens with increasing N2O flow rate, consistent with above-mentioned Raman (Zni) and XPS (VO and Zni) results. The g ~2.002 signal has been associated with a hole residing in one of the non-axial oxygen atoms at the neighboring site of a zinc vacancy [44]. The emergence of this resonance for sample D indicates the easier formation of VZn related acceptors when the growth takes place in O-rich ambience.
Fig. 4 EPR spectra of samples A and D. The inset shows the schematic diagram of the measurement setup.
3.5. PL characterization
Optical characterization of the N-doped ZnO MRs was carried out in order to gain more information on the intrinsic point defect properties. Firstly, we took the PL spectra of sample D as an example to identify the emission lines. TD-PL spectra of sample D were recorded in the temperature range of 13 - 160 K. The near band edge (NBE) emissions from 3.21 to 3.39 eV and the energy positions of these emissions as a function of temperature are illustrated in Figs. 5(a) and 5(b). The highest energy peak at 3.377 eV is free excitons (FX) according to the literature [45]. With the relation
the continuous redshift of FX energy position EFX(T) with increasing temperature owing to the bandgap shrinkage could be fitted by the Varshni’s equationwhere Eg(0) is the bandgap at 0 K, while α and β are constants. As shown by the solid fitting line in Fig. 5(b), we can get α = 7.82 × 10−4 eV/K and β = 831 K, close to the values quoted in Ref. [45]. Considering the large surface to volume ratio of ZnO MRs, the emission at 3.368 eV is assigned to surface bound excitons (SX) with its intensity decreasing dramatically fast with increasing temperature. The emissions at 3.363, 3.359, and 3.240 eV are assigned to D0X, A0X and DAP recombination, respectively. The D0 and A0 here are in association with the Zni related shallow donors and the VZn related shallow acceptors, respectively, as has been discussed previously [25]. Furthermore, as the 0.5kBT offset between solid and dashed lines clearly shown in Fig. 5(b), which can be used as an evidence to judge whether an emission is eA0 or not, we ascribe the peak at 3.311 eV to the radiative recombination of free electron to shallow acceptor, with its energy position described aswhere EA is the acceptor binding energy and calculated to be ~126 meV. Lastly, the two peaks at 3.289 and 3.218 eV are ascribed to the first and second longitudinal optical (LO) phonon replica of A0X because of the energy spacing, and denoted as A0X-LO and A0X-2LO.Fig. 5 TD-PL spectra of sample D in the temperature range of 13 - 160 K. (a) NBE emissions. (b) The energy positions of the NBE peaks as a function of temperature. (c) GB emissions. The inset is the integrated intensity of GB emissions normalized to the value at 13 K as a function of temperature, where the solid line is fitting curve of the experimental results by Eq. (4).
The green band (GB) emissions of sample D with zero-phonon-line (ZPL) around 2.85 eV and its 8 LO phonon replicas are presented in Fig. 5(c). As plotted in the inset, the integrated intensity of GB emissions increases with increasing temperature until 40 K, which is also known as a negative thermal quenching (NTQ) effect [46,47]. The solid line shows an excellent agreement between experiment data and the nonlinear least-square fitting with the equation
where C1, C2 and C3 are constants, E1 is the activation energy of NTQ, while E2 and E3 are the activation energies of the nonradiative recombination, representing the donor binding energy of the shallow donor and the acceptor binding energy of the deep acceptor [9], which have been calculated to be 26.6 and 436 meV, respectively. This result verifies that the origin of the GB emissions with fine structure is the radiative transitions from the ground and exited states of the shallow donor Zni recombining with deep acceptor isolated VZn [25].Finally, to analyze the differences of the optical properties between the samples grown at different flow rates of N2O, PL spectra measured at 9 K are depicted in Fig. 6(a) with a resolution of 0.05 nm. In the visible range, all the GB emissions of the four samples exhibit fine structures, which consist of a doublet with a fixed splitting energy of 27 meV, repeated with a LO phonon energy spacing of 72 meV. Besides, at the lower energy side in the UV region, the PL spectra are all dominated by the broad DAP recombination around 3.240 eV and its 4 LO phonon replicas. Obviously, the intensity of GB emission and DAP recombination increases from sample A to B, then decreases from sample B to D, showing the same trend with the AMs in Raman spectra. It implies that the donor responsible for DAP and GB emissions is related to Zni. Consequently, we conclude that the shallow donor defect Zni in N-doped ZnO MRs can be suppressed at higher flow rates of N2O. Furthermore, from the NBE emissions shown in Fig. 6(b), we have observed four peaks with the same energy positions at 3.377, 3.368, 3.363, and 3.359 eV, which have been identified as FX, SX, D0X, and A0X, respectively. It is encouraging to see that as the flow rate of N2O increases, the intensity of A0X increases monotonically and becomes the main peak for samples C and D, combined with the appearance of A0X-LO and A0X-2LO as shown in Fig. 6(a). Considering the EPR results, the trend implies that the shallow acceptors in our N-dope ZnO MRs, possibly related to VZn or VZn clusters [23,25], can be enhanced with higher flow rate of N2O.
Fig. 6 (a) PL spectra of the MR samples measured at 9 K. (b) The NBE emissions from 3.34 to 3.39 eV.
4. Conclusion
In summary, vertically aligned N-doped ZnO MRs with excellent crystallinity were carried out via the chemical vapor transport method at different flow rates of N2O, which was employed as both N and O precursors. With the flow rates of N2O increasing, the shallow donor Zni related Raman modes, optical emissions, and EPR signals decrease, while the shallow acceptor VZn related EPR signals and optical emissions increase. Our results show that the native defects can be controlled by simply changing the flow rate of N2O during the growth of ZnO MRs by CVT method. More work is ongoing to optimize the experiment conditions and finally achieve reliable p-type conductivity in N-doped ZnO MRs.
Acknowledgments
This research was supported by the State Key Program for Basic Research of China under Grant No. 2011CB302003, National Natural Science Foundation of China (Nos. 61274058, 61322403, 61504057, and 61574075), the Natural Science Foundation of Jiangsu Province (Nos. BK20130013 and BK20150585), the Six Talent Peaks Project in Jiangsu Province (2014XXRJ001).
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