- Email: [email protected]
GaN nanophosphors for white-light applications
Optical Materials 75 (2018) 61e67
Contents lists available at ScienceDirect
Optical Materials journal homepage: www.elsevier.com/locate/optmat
GaN nanophosphors for white-light applications Mirgender Kumar a, V.P. Singh b, Sarvesh Dubey c, Youngsuk Suh a, **, Si-Hyun Park a, * a
Department of Electronics Engineering, Yeungnam University, Gyeongsan 38541, South Korea School of Engineering, Indian Institute of Technology, Mandi 175005, India c Department of Physics, L. N. D. College, Muzaffarpur 845401, India b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 June 2017 Received in revised form 13 October 2017 Accepted 13 October 2017
GaN nanoparticles (NPs) were synthesized by carbothermal reduction combined with nitridation, using Ga2O3 powder and graphitic carbon nitride (g-C3N4) as precursors. Characterization of the NPs was performed by X-ray diffraction, scanning electron microscopy, and room-temperature photoluminescence measurements. X-ray photoelectron spectroscopy was also performed to detect the chemical states of the different species. A universal yellow luminescence (YL) band was observed from complexes of Ga vacancies with O anti-sites and of O anti-sites with C. Further increments in the C content were observed with continued growth and induced an additional blue luminescence (BL) band. Tuning of the YL and BL bands resulted in white-light emission under certain experimental conditions, thus offering a new way of employing GaN nanophosphors for solid-state white lighting. Calculations of the correlated color temperature and color-quality scale parameters confirmed the utility of the experimental process for different applications. © 2017 Elsevier B.V. All rights reserved.
Keywords: GaN nanoparticles White-light emission Correlated color temperature Color-quality scale
1. Introduction GaN shows great potential in electronics and optoelectronics because it has a wide band gap, good thermal properties, high luminescence, strong interatomic bonds, and high carrier mobility, suitable for application in light-emitting diodes (LEDs), lasers, and high-electron-mobility transistors (HEMTs) [1,2]. Ternary alloys of GaN have also reached a significant plateau in semiconductor technology and are used over a large spectral range from infrared (IR) to deep ultraviolet (UV) electromagnetic radiation [3,4]. In addition, the nontoxic nature of GaN offers an advantage over Asand phosphide-based semiconductors, which have been studied extensively in recent decades. Therefore, extensive efforts have been devoted to studying GaN in the forms of thin films and nanostructures with different morphologies, fabricated using different techniques to explore new phenomena for novel applications [5e7]. Many breakthroughs in thin-film growth have been
Abbreviations: BL, blue luminescence; CCT, correlated color temperature; CQS, color-quality scale; SEM, scanning electron microscopy; XPS, X-ray photoelectron spectroscopy; XRD, X-ray diffraction; YL, yellow luminescence. * Corresponding author.. ** Corresponding author. E-mail addresses: [email protected] (Y. Suh), [email protected] (S.-H. Park). https://doi.org/10.1016/j.optmat.2017.10.021 0925-3467/© 2017 Elsevier B.V. All rights reserved.
reported, especially for the UV/blue region; however, these remain non-cost effective for device applications. Therefore, nanoscale structures of undoped GaN may be good candidates for different applications and thus merit further research. Most GaN nanostructures have been studied, such as nanowires, nanorods, nanobelts, nanopillars, nanoparticles (NPs), and quantum dots, including special coreeshell structures [8e12]. Undoped GaN NPs have gained much attention because they have emission properties in the green, yellow, blue, and red bands of the visible spectrum, which enhance their applicability in a wide range for solid-state lighting and sensors [13e16]. These unprecedented properties have evolved from the inherent surface states accompanying vacancies (Ga vacancy, VGa, and N vacancy, VN), which act as luminescence centers and contribute to visible emission. Some surface states are generated as dangling bonds, which may act as trap centers rather than contributing as luminescence center. Some vacancy-type surface states may be occupied by residual impurities as well as anti-site atoms, which also contribute to different luminescence bands of the energy spectrum. Further, these surface states (donors or acceptors) can also form complexes with opposite-polarity states. Thus, NP surface states and complexes thereof can induce defect energy levels found between the band edges of corresponding energies and correlated well to their respective bands of luminescence [17e20]. A broad yellow luminescence (YL) band centered at 2.1e2.3 eV is a widely studied
62
M. Kumar et al. / Optical Materials 75 (2018) 61e67
visible band in the photoluminescence (PL) of GaN. An early study established that radiative transition from a shallow donor to a deep acceptor causes this YL band. In addition, by examining the broad and Gaussian nature of the emission spectrum, many studies have associated the origin of the YL band with the transition from the deep donor to shallow acceptor state, along with two transition involvement models from the shallow donor to the deep acceptor and from the deep donor to the shallow acceptor, which involve the complexes of VGa surface states and complexes of carbon defects at VN surface states with the nearest neighbors [21e24]. Some studies have also reported high-purity GaN showing a broad PL feature in the visible range, explained by de-convolution into two bands of YL centered at 2.2 eV and green luminescence (GL) centered at 2.48 eV [25,26]. The shifting of the YL band into the green region has also been observed with increases in excitation intensity, signifying that both bands correspond to the same defects with different charge or complexation states [27]. The blue luminescence (BL) band in the visible spectrum has also been studied significantly, with unstable properties assigned to specific defects [28,29]. The contribution of VGa surface states with some residual impurities has been stated as a probable reason for this band. High growth temperatures, which induce Ga desorption in N-deficient ambient atmosphere, were found to favor the generation of many VGa surface states. In addition to this correspondence, the combination of shallow donors (C occupying VGa surface states, CGa) with shallow acceptor levels (C occupying VN surface states, CN) has also been reported as a convincing cause. Some of the studies have also correlated the BL and YL band intensities with intentional or unintentional C doping, which may induce interstitial (CI) and anti-site (CN and CGa) defects [30,31]. These defects can also originate from intrinsic surface states relating to N and Ga vacancies, which can be tuned effectively in NPs. Some recent studies have reported white-light emission from GaN superlattice structures by tuning the most prominent YL and BL bands arising from interfacial state-related defects and residual impurities [32,33]. In this regard, GaN NPs might have much potential because they show excellent emission tuning capabilities by the modification of surface morphology through different synthesis processes. However, no process yet reported has successfully tuned the YL and BL bands for GaN NPs, despite their attributes as strong contenders for next-generation solid-state tunable-colortemperature white-light applications. Among all of the techniques available for synthesizing NPs, nitridation is an effective route regarding cost effectiveness, controllability of operating conditions, single-phase production, and high product purity. Generally, these reactions are performed using Ga precursors with N2 and NH3 gases at high temperatures. The low operating temperature and ease of use has made the nitridation of Ga2O3 with C and NH3 the preferred route, known as the carbothermal nitridation process [34e38]. An NH3-free nitridation process has also emerged as an alternative method with graphitic carbon nitride (gC3N4), which can be performed at a much lower temperature with better surface morphology [39,40]. Therefore, this work investigates the synthesis of GaN NPs for white-light applications by carbothermal reduction and nitridation, using Ga2O3 powder and g-C3N4 as the source materials. The synthesis was performed under different experimental conditions to tune the emission properties over a broad visible range with the variation in the concentration of residual O and C impurities. GaN NPs with tuned emission properties are good candidates for the further development of variable-color-temperature nanophosphors for white light-emitting diodes (LEDs) with good efficiency.
2. Experimental process Ga2O3 powder and g-C3N4 were used as the Ga precursor and nitridation reagent, respectively. The g-C3N4 was prepared by the thermal condensation of dicyandiamide, which was used to permit a low-temperature nitridation reaction with high-crystallinity products having good surface morphology. Ga2O3 and dicyandiamide were purchased from Sigma Aldrich with 99% purity. g-C3N4 is widely used and decomposes at 600 C into highly reactive C and N. C has strong reducing capabilities toward metal oxides, and nitrides can form in the presence of nitride ions. Hence, g-C3N4 carries the capability of reducing metal oxides into their corresponding metals followed by nitridation thereof, which solves the problem of controlling the rates of reduction and nitridation. Another advantage of using g-C3N4 is the formation of highly stable CO and N2 as byproducts. In comparison, NH3-based nitridation processes involve H2 in the reaction and degrade the morphology of the final product. A typical nitridation reaction was performed in a tubular furnace by keeping the grounded mixture of Ga2O3 and g-C3N4 (1:3 M ratio) in an alumina crucible under flowing N2 (5 mL/min) at 800 C (sample S1) or 750 C (sample S2) at atmospheric pressure. Further, samples S3 and S4 were synthesized at 750 C with the pressures of 50 Torr (6.7 kPa) and 100 Torr (13 kPa), respectively. Sample S5 was prepared with the reaction furnace pressure of 100 Torr (13 kPa) and the temperature of 650 C. Finally, samples of the GaN NPs were characterized by X-ray diffraction (XRD), scanning electron microscopy (SEM), PL spectroscopy, and X-ray photoelectron spectroscopy (XPS). 3. Results and discussion Fig. 1 shows the XRD pattern obtained using a Philips X'pert Pro
Fig. 1. XRD pattern of all samples for GaN NPs.
M. Kumar et al. / Optical Materials 75 (2018) 61e67
diffractometer with Cu Ka radiation for GaN samples using the carbothermal nitridation process with different synthesis parameters. The analysis of the XRD patterns reveals the formation of hexagonal wurtzite-structured GaN with different features corresponding to the (100), (002), (101), (102), (110), (103), (112), and (201) crystallographic planes. The structural analyses of all samples are summarized in Table 1 corresponding to the (101) dominating features. The average particle size, calculated using the Scherrer formula, for sample S2 is found to decrease because of the decrease in reaction temperature, which decreases the growth rate of NPs. However, the average particle size is increased for samples S3 and S4 because of the increase in growth rate with the decrease in ambient pressure. Similar to S2, the average particle size for S5 is also decreased compared to that of S4 with the decrease in reaction temperature. No XRD features corresponding to Ga2O3 and C3N4 are observed, confirming the full conversion of Ga2O3 to GaN. SEM image (Fig. 2(a)) using EVO-40 ZEISS depicts the surface morphology of the GaN NPs for sample S5. The reaction product comprises NPs with platelet shapes. Fig. 2(b) shows the histogram of NP size distribution for sample S5. Energy-dispersive X-ray (EDX) spectra of different samples, shown in Fig. 3, were obtained for the qualitative elemental analysis of different elements in the GaN samples synthesized under various experimental conditions. The EDX spectrum of sample S1 shows a small signal corresponding to C and O along with another intense signal for Ga and N, possibly originating from the precursors used in the reaction. Therefore, the XRD pattern and EDX elemental signals of C and O confirm their presences as residual impurities. The signal corresponding to C and O is stronger for sample S2 compared to that for S1 because of the lower reaction temperature. The same trends for C- and O-related signals are observed for samples S3 and S4, which are prepared at lower pressures than those for S1 and S2. Thus, impurities related to C and O are enhanced with decreases in reaction temperature and pressure. Quantitative elemental analyses are summarized in Table 2. Fig. 4 shows the PL spectra of the GaN NPs samples (S1, S2, S3, S4, and S5) obtained with the excitation wavelength of 325 nm from a HeeCd laser and using a photomultiplier tube as a detector. Sample S1 exhibits two different luminescence bands in the visible and UV ranges. The narrow UV luminescence band (~3.4 eV) is a typical signature of GaN, corresponding to near-band-edge emission. The broad visible band centered at 2.25 eV confirms the presence of defects within the band gap, induced by surface states (VGa and VN) and their anti-site-including complexes. This visible band is recognized as the well-known yellow band, reported by many researchers. The breadth of the band is attributed to the combination of multiple transitions, possibly including transitions from shallow donors to deep acceptors and from deep donors to shallow acceptors in this case. According to the impurities detected by EDX, O anti-site defects (O at N vacancy-type surface states, ON) and C antisite defects (C at N vacancy-type surface states, CN) probably serve as the shallow donors and acceptors, respectively, residing below the conduction band edge (CBE) at 0.30 eV and above the valence band edge (VBE) at 0.23 eV, respectively. The deep acceptor and donor may be complexes of different isolated defect levels,
Table 1 Structural parameters of the GaN NPs corresponding to (101) crystallographic plane. Sample
Position of peak 2q (deg.)
FWHM (deg.)
Average particle size (nm)
S1 S2 S3 S4 S5
36.92 36.90 36.90 36.87 36.85
0.8750 0.9487 0.7543 0.3928 0.5434
95.7 88.2 111 212.9 154.1
63
recognized as a complex of Ga vacancy-type surface states (VGa) and O anti-site defects (VGe(ON)n) (n ¼ 1e3) and a complex of C antisite defects with O anti-site defects (CNeON). VGa is an intrinsic defect in GaN and its presence is highly probable in GaN NPs; it may act as a deep acceptor at ~1.1 eV above the VBE. The VGa defects mostly exist in complexes with donors such as ON, rather than as isolated states. The conditions of VG as acceptor and ON as donor also favor the formation of (VGe(ON)n) complexes. Furthermore, the much smaller distinction between the YL and GL may also indicate the involvement of the same type of closely spaced levels, such as (VGaeON) for YL and (VGae(ON)2) for GL. The shifting of VGa closer to the VBE by forming a complex with ON may also indicate the presence of multiple types of complex, which would explain the greenish-yellow luminescence. A similar scenario also exists for the CNeON complex defect because of the different polarities of the constituent defects. However, their energy positions remain unconfirmed and require further investigation. Therefore, the VGainduced surface states and VN-induced surface states occupied by O and C are mainly responsible for the YL in the synthesized GaN NPs. The yellow emission band is stronger in the spectrum of sample S2, as well as broader and slightly shifted to a higher energy, because of the lower reaction temperature, at the expense of a weakened UV band. The formation of defects in sample S2 is clearly promoted by the decrease in temperature, which causes an increase in the intensity of the YL band. This phenomenon may occur with the decreased average size of the NPs as observed by structural analysis, which enhances the surface-to-volume ratio and creates higher numbers of surface induced impurities states. The stronger C and O signals in the EDX spectrum of sample S2 compared to those of S1 also indicate higher defect density in S2. Hence, the higher level of O in the sample may also initiate the formation of defect levels belonging to (VGae(ON)2) and (VGaeON), resulting in a shift of the luminescence peak toward a higher energy, as observed here. In response, the intensity of the UV band is quenched because of the weaker free exciton recombination with the involvement of different radiative and non-radiative defect levels, related to the enhanced surface states with the reduced surface-to-volume ratio. A new BL band appears in the spectrum of sample S3 along with a stronger YL band, as observed in the spectra of S1 and S2. The BL band intensity is lower than that of the YL band, however, and the stronger visible bands suppress the UV band. The BL band in the visible region may arise from the transition from a shallow donor defect level to a shallow acceptor defect level [41]. The existence of both BL and YL bands may also indicate the formation of a new shallow donor along with ON, expected to be near the CBE. In this case, the most probable donor state is the C anti-site defect at the Ga site (CGa) at 0.20 eV below the CBE [17]. However, this defect level requires higher formation energy. The increase in C impurities with the decrease in reaction pressure, along with the abundance of Ga vacancies, may favor the generation of CGa. Therefore, the increased level of residual impurities probably causes the strengthening of the YL band and the inception of the BL band. A further decrease in the operating pressure of the reaction to prepare sample S4 increases the intensities of both YL and BL. The BL band intensity is increased significantly compared to that of the YL band because of the larger role of CGa as a shallow donor along with the generation of interstitial O (OI) and C (CI). The presences of interstitial O and C have been confirmed experimentally in GaN, and reportedly contribute to the visible bands [42,43]. The YL band continues to increase in intensity probably because of the increasing amount of C impurities, which also induce the generation of a new mid-gap defect level corresponding to interstitial C. Further, the PL spectrum of sample S5 shows the increased intensity of both visible bands, which appear almost overlapped upon the decrease in reaction temperature, compared to those of sample S4
64
M. Kumar et al. / Optical Materials 75 (2018) 61e67
Fig. 2. SEM image (a), and the histogram of size distribution (b) for GaN NPs sample S5.
Fig. 3. EDX spectra of different NPs samples prepared with different experimental conditions.
Table 2 Percentage of different elements in GaN NP samples calculated from EDX spectra. Sample/Elements
Ga (%)
N (%)
C (%)
O (%)
S1 S2 S3 S4
81 77 72 68
14 10 8 7
2 5 9 12
3 8 11 13
with the same reaction chamber pressure. This is expected because of the increase in residual impurity levels, which enhance the generation of contributing defect levels. Moreover, the transition energy between ON and OI (~0.42 eV above the VBE) may also contribute to this broad BL, which is highly feasible with a high level of residual O [42]. The broad luminescence obtained in the white-light region is consistent with the above discussion and with previously reported results. Thus, the process of variable experimental conditions establishes GaN NPs as strong candidates for UV-excited phosphors in solid-state lighting, as they generate broad visible emission with controlled C and O contents.
Fig. 4. PL spectra with the discussed different experimental conditions.
XPS measurements were also performed for synthesized GaN NPs of sample S5 to examine the chemical state of different elements, which can confirm the existence of different defect levels. Fig. 5 shows the different XPS peaks corresponding to Ga 3d, N 1s, O 1s, and C 1s. The Ga 3d peak in Fig. 5(a) shows a slight asymmetry on the higher binding energy side. By fitting this spectrum with a Gaussian profile, it can be de-convoluted into two components, G1 (~19.7 eV) and G2 (~20.8 eV). G1 represents the dominant peak corresponding to the GaeN bond, and the higher-binding-energy peak of G2 corresponds to the GaeO bond. G2 is observed at a higher binding energy with a greater difference in the electronegativity of Ga and O compared to that of Ga and N [44]. This higherbinding-energy peak is usually attributed to oxidation states, oxynitride, GaeOH components, and hybridization states of N 2s and Ga d states. The possibility of GaeOH is negligible because H has no role in the reaction. The options of oxide and oxynitride states are also minimal, with no supporting XRD peaks. Hybridization state features also have very low possibilities in the presence of many residual O impurities. However, the present reaction conditions of low ambient pressure and temperature along with a high concentration of residual O is expected to favor O anti-site (ON) defect formation instead of the above discussed cases. This is also supported by the EDX
M. Kumar et al. / Optical Materials 75 (2018) 61e67
65
Fig. 5. XPS for synthesized GaN NPs: (a) Ga 3d peak; (b) N 1s peak; (c) O 1s peak; (d) C 1s peak.
measurements, which show a much lower concentration of N than that for Ga and suggests the partial incorporation of O into the GaN lattice by substitution of N vacancies. Finally, our reaction was performed at a lower pressure, which favors the fast carbothermal reduction of Ga2O3 and therefore removes the possibility of oxide states. If Ga2O3 were present, then it would appear in XRD because of the crystalline nature of the precursor used in the reaction. Therefore, the G2 signature is attributed to the presence of O at N sites, confirming the existence of O anti-site defects (ON) [45]. The fitted N 1s spectrum in Fig. 5(b) shows two de-convoluted peaks, N1 (~397.2 eV) and N2 (~398.2 eV). The N1 peak belongs to the NeGa bond, and the peak N2 at higher binding energy can be attributed to the CeN bond because of the higher electronegativity difference of C and N compared to that of N and Ga. Consequently, the N2 peak indicates the presence of elemental C at Ga sites, thus confirming the existence of C anti-site defects at Ga sites (CGa). Fig. 5(c) shows the O 1s spectrum of GaN NPs de-convoluted into two peaks, O1 (~530.4 eV) corresponding to the OeGa bond, and O2 (~531.1 eV) for interstitial O (OI). Hence, O1 represents the same defects state as G2 for ON, and peak O2 represents OI defects as acceptors close to the VBE [42]. After Gaussian fitting, the spectrum of C 1s shown in Fig. 5(d) is de-convoluted into three peaks, C1 (~287 eV), C2 (~286 eV), and C3 (~285 eV). The peak with lowest binding energy, C3, represents elemental C, which may originate from residual C. In addition, the diffusion of C is highly likely; C may behave as a deep acceptor in the form of an interstitial atom (CI) [43]. The high C concentration found in the EDX measurements as a residual impurity also supports the formation of interstitial states
under a low reaction chamber pressure. The higher-energy peak C2 represents the CeN bond and indicates C substitution at the Ga sites (CGa), which are also detected in the N 1s spectrum. Another higher-binding-energy peak, C1, is also shifted because of the higher difference in electronegativity between C and Ga in the CeGa bond, i.e., C substitution at N sites (CN). Fig. 6 shows the Commission Internationale de l'Eclairage (CIE) chromaticity coordinates obtained from the PL spectra of the GaN samples. The inset shows the corresponding correlated color temperature (CCT) along with the color quality scale (CQS) parameter. The coordinates of samples S1 and S2 are located in the yellow region in the chromaticity diagram, as expected. However, the coordinates of S3 are shifted to the white region because of the inception of the blue band. The coordinates of samples S4 and S5 are closer to the standard white light coordinates (0.33, 0.33) because of the reduced intensity difference between the YL and BL bands. The value of CCT for samples S1 to S5 increase continuously from 3800 to 6000 K, while the CQS value of sample S2 is lower than that of S1 and thereafter increases consistently from 89 for S3 to 95 for S5. The quality of white light of S5 appears better than those of S3 and S4 with an increased CQS, showing luminescence equivalent to that of cool daylight. Fig. 7 shows the energy level diagram of the synthesized GaN NPs, explaining the contributions of different defect levels to the different visible bands that attain broad white luminescence. As shallow donors, ON and CGa have energy levels at 0.30 and 0.20 eV below the CBE, respectively [17,46]. Indeed, the high defect concentration and lattice interaction may also induce broadening of an
66
M. Kumar et al. / Optical Materials 75 (2018) 61e67
acceptor (CN) confirms the contribution to the yellow band, which also supports the hypothesis of Glaser et al., who suggested a twostep YL transition with non-radiative capture from a shallow donor to a deep donor and radiative recombination from a deep donor to a shallow acceptor [23]. Julkarnain et al. confirmed a deep acceptor at 1 eV for the contribution to YL [48], with an earlier hypothesis of transition from a shallow donor to a deep acceptor. In our case, interstitial C is considered another deep acceptor, which evolves with high levels of residual C impurities in the samples under favorable experimental conditions. Guzm an et al. also confirmed that interstitial O at 0.42 eV above the VBE contributes to BL [42]. The intensification of both visible bands is consistent with the increasing defect states related to C and O. Our proposed band diagram also explains the YL band for GaN samples with and without C impurities. Therefore, the visible BL, GL, and YL can be converted into white-light emission. 4. Conclusions
Fig. 6. Chromaticity coordinates (CIE 1931) for the different samples, where C ¼ S1, ◄ ¼ S2, A ¼ S3, ; ¼ S4. Inset: corresponding CCT and CQS.
GaN NPs were successfully synthesized by carbothermal reduction and nitridation, as confirmed by XRD, EDX, and SEM analysis. White luminescence was also attained by varying the experimental conditions, which generated different defect levels between the band edges. The ability to tune defects related to C and O demonstrated that GaN NPs were promising candidates for nanophosphor-based white-light sources. The proposed experimental process also demonstrated the tunability of color temperature along with color quality scale for different lighting purposes. These results may enhance the further development of GaN-based devices with improved performance. Acknowledgement This work was supported by the 2014 Yeungnam University Research Grant. References
Fig. 7. Schematic of the proposed energy level diagram of the GaN NPs responsible for perceived white-light emission.
individual defect state. Therefore, the two closely spaced shallow donors can be considered localized states between 0.15 and 0.30 eV, suggesting non-radiative relaxation from the CBE and thereby supporting broad emission for different visible bands. As discussed above, the three complexes of Ga vacancies and O anti-site defects may exist in conjunction with O anti-site defects (n ¼ 1e3). However, only VGe(ON) and VGe(ON)2 can promote the emission of YL and GL, respectively. The VGa level is located 1.1 eV above the VBE, and moves to 0.9 eV after forming the complex VGe(ON) [45]. Therefore, VGe(ON)2 should shift to 0.8 eV, closer yet to the VBE. In addition, the O-containing samples exhibit higher probabilities of forming VGa and ON complexes than that for forming an isolated VGa [14]. The hypothesis of the formation of VGe(ON)2 may be supported by these facts. The possibility of a deep donor (CNeON) contributing to YL supports the scenario of two transitions creating YL, as proposed by Polyakov et al. [47]. The energy state at 0.9 eV (below the CBE) of this deep donor and the transition to a shallow
[1] T.J. Flack, B.N. Pushpakaran, S.B. Bayne, GaN technology for power electronic applications: a review, J. Electron. Mater. 45 (2016) 2673. [2] H. Morkoç, Handbook of Nitride Semiconductors and Devices, GaN-based Optical and Electronic Devices, John Wiley & Sons, 2009. [3] S. Nakamura, T. Mukai, M. Senoh, Candela-class high-brightness InGaN/AlGaN double-heterostructure blue-light-emitting diodes, Appl. Phys. Lett. 64 (1994) 1687. [4] A.G. Bhuiyan, K. Sugita, A. Hashimoto, A. Yamamoto, InGaN solar cells: present state of the art and important challenges, IEEE J. Photonics 2 (2012) 276e293. [5] E.D.L. Boulbar, C.J. Lewins, D.W.E. Allsopp, C.R. Bowen, P.A. Shields, Fabrication of high-aspect ratio GaN nanostructures for advanced photonic devices, Microelectron. Eng. 153 (2016) 132e136. [6] M.M. Rozhavskaya, W.V. Lundin, E.E. Zavarin, E.Y. Lundina, S.I. Troshkov, V.Y. Davydov, M.A. Yagovkina, P.N. Brunkov, A.F. Tsatsulnikov, Various types of GaN/InGaN nanostructures grown by MOCVD on Si(111) substrate, Phys. Status Solidi C 10 (2013) 441e444. [7] M. Nyka, W. Strek, J.M. Jabonski, L. Kepinski, R. Kudrawieca, J. Misiewicz, Synthesis, structure and optical properties of GaN nanocrystallites, Mater. Sci. Semicond. Process. 8 (2005) 511. [8] P. Corfdir, C. Hauswald, O. Marquardt, T. Flissikowski, J.K. Zettler, S.F. Garrido, L. Geelhaar, H.T. Grahn, O. Brandt, Crystal-phase quantum dots in GaN quantum wires, Phys. Rev. B 93 (2016) 115305. [9] S. Takeda, K. Ishino, Y. Inoue, A. Ishida, H. Fujiyasu, H. Kominami, H. Mimura, Y. Nakanishi, S. Sakakibara, Growth of pillarlike GaN nanostructures, Jpn. J. Appl. Phys. 44 (2005) 5664. [10] Z. Zang, X. Zeng, J. Du, M. Wang, X. Tang, Femtosecond laser direct writing of microholes on roughened ZnO for output power enhancement of InGaN lightemitting diodes, Opt. Lett. 41 (2016) 3463. [11] L. Zhang, C. Liu, White light emission from GaN stack layers doped by different rare-earth metals, Jpn. J. Appl. Phys. 54 (2015) 02BA03. [12] M. Nyk, R. Kudrawiec, W. Strek, J. Misiewicz, Synthesis and optical properties of Eu3þ and Tb3þ doped GaN nanocrystallite powders, Opt. Mater. 28 (2006) 767. [13] Y.H. Ko, J. Song, B. Leung, J. Han, Y.H. Cho, Multi-color broadband visible light source via GaN hexagonal annular structure, Sci. Rep. 4 (2014) 5514.
M. Kumar et al. / Optical Materials 75 (2018) 61e67 [14] X. Li, P.W. Bohn, J.J. Coleman, Impurity states are the origin of yellow-band emission in GaN structures produced by epitaxial lateral overgrowth, Appl. Phys. Lett. 75 (1999) 4049. [15] C. Li, Z. Zang, W. Chen, Z. Hu, X. Tang, W. Hu, K. Sun, X. Liu, W. Chen, Highly pure green light emission of perovskite CsPbBr3 quantum dots and their application for green light-emitting diodes, Opt. Express 24 (2016) 15071. [16] X. Shi, M. Peng, J. Kou, C. Liu, R. Wang, Y. Liu, J Zhai, A Flexible GaN nanowire array-based schottky-type visible light sensor with strain-enhanced photoresponsivity, Adv. Electron. Mater. 1 (2015) 1500169. [17] A.F. Wright, Substitutional and interstitial carbon in wurtzite GaN, J. Appl. Phys. 92 (2012) 2575. [18] A. Sedhain, J. Li, J.Y. Lin, H.X. Jianga, Nature of deep center emissions in GaN, Appl. Phys. Lett. 96 (2010) 151902. [19] N.T. Son, C.G. Hemmingsson, T. Paskova, K.R. Evans, A. Usui, N. Morishita, n, Identification of the gallium T. Ohshima, J. Isoya, B. Monemar, E. Janze vacancyeoxygen pair defect in GaN, Phys. Rev. B 80 (2009) 153202. [20] B. Schwenzer, S. Keller, L. Loeffler, R. Seshadri, F.F. Lange, S.P. Baars, U.K. Mishra, Group-III nitride nanoparticles-synthesis and photoluminescence studies, in: Conference Proceeding of IEEE-NANO, 2003, pp. 12e14. [21] T. Ogino, M. Aoki, Jpn. Mechanism of yellow luminescence in GaN, J. Appl. Phys. 19 (1980) 2395. [22] D.M. Hofmann, D. Kovalev, G. Steude, B.K. Meyer, A. Hoffmann, L. Eckey, R. Heitz, T. Detchprom, H. Amano, I. Akasaki, Properties of the yellow luminescence in undoped GaN epitaxial layers, Phys. Rev. B 52 (1995) 16702e16706. [23] E.R. Glaser, T.A. Kennedy, K. Doverspike, L.B. Rowland, D.K. Gaskill, J.A. Freitas, M.A. Khan, D.T. Olson, J.N. Kuznia, D.K. Wickenden, Optically detected magnetic resonance of GaN films grown by organometallic chemical-vapor deposition, Phys. Rev. B 51 (1995) 13326e13336. [24] E. Calleja, F.J. Sanchez, D. Basak, M.A.S. Garcia, E. Munoz, I. Izpura, F. Calle, J.M.G. Tijero, J.L.S. Rojas, B. Beaumont, P. Lorenzini, P. Gibart, Yellow luminescence and related deep states in undoped GaN, Phys. Rev. B 55 (1997) 4689. [25] M.A. Reshchikov, H. Morkoc, S.S. Park, K.Y. Lee, Yellow and green luminescence in a freestanding GaN template, Appl. Phys. Lett. 78 (2001) 3041e3043. [26] M.A. Reshchikov, D.O. Demchenko, A. Usikov, H. Helava, Y. Makarov, Carbon defects as sources of the green and yellow luminescence bands in undoped GaN, Phys. Rev. B 90 (2014) 235203e235211. [27] M.A. Reshchikov, H. Morkoc, S.S. Park, K.Y. Lee, Two charge states of dominant acceptor in unintentionally doped GaN: evidence from photoluminescence study, Appl. Phys. Lett. 81 (2002) 4970e4972. [28] M. Gopalakrishnan, V. Purushothaman, V. Ramakrishnan, G.M. Bhaleraoc, K. Jeganathan, The effect of nitridation temperature on the structural, optical and electrical properties of GaN nanoparticles, CrystEngComm 16 (2014) 3584e3591. [29] Y. Yang, V.J. Leppert, S.H. Risbud, B. Twamley, P.P. Power, H.W.H. Lee, Blue luminescence from amorphous GaN nanoparticles synthesized in situ in a polymer, Appl. Phys. Lett. 74 (1999) 2262e2264. ndez, R.M. Pe rez, B.M. Monroy, A.E. Esquivel, [30] G. Santana, O. Melo, J.A. Herna pez, F. Moure, L.A. Hern M.L. Lo andez, G.C. Puente, Photoluminescence study of gallium nitride thin films obtained by infrared close space vapor transport,
67
Materials 6 (2013) 1050e1060. [31] J. Kang, Y. Shen, Z. Wang, Effects of residual C and O impurities on photoluminescence in undoped GaN epilayers, Mater. Sci. Eng. B91e92 (2002) 303e307. [32] N.B. Sedrine, T.C. Esteves, J. Rodrigues, L. Rino, M.R. Correia1, M.C. Sequeira, A.J. Neves, E. Alves, M. Bockowski, P.R. Edwards, K.P. O'Donnell, K. Lorenz, T. Monteiro, Photoluminescence studies of a perceived white light emission from a monolithic InGaN/GaN quantum well structure, Sci. Rep. 5 (2015) 13739e13741. [33] M.L. Lee, Y.H. Yeh, S.J. Tu, P.C. Chen, W.C. Lai, J.K. Sheu, White emission from non-planar InGaN/GaN MQW LEDs grown on GaN template with truncated hexagonal pyramids, Opt. Express 23 (2015) A401eA412. [34] C.M. Balkas¸, R.F. Davis, Synthesis routes and characterization of high-purity, single-phase gallium nitride powders, J. Am. Ceram. Soc. 79 (1996) 2309e2312. [35] T.H. Kim, S. Choi, D.W. Park, Thermal plasma synthesis of crystalline gallium nitride nanopowder from gallium nitrate hydrate and melamine, Nanomaterials 6 (2016) 38. [36] V. Ganesh, S. Suresh, M. Balaji, K. Baskar, Synthesis and characterization of nanocrystalline gallium nitride by nitridation of Ga-EDTA complex, J. Alloys Compd. 498 (2010) 52e56. [37] W.S. Jung, Reaction intermediate(s) in the conversion of h-gallium oxide to gallium nitride under a flow of ammonia, Mater. Lett. 57 (2002) 110e114. [38] W.S. Jung, Reaction mechanism for the ammonolysis of b-gallium oxide to gallium nitride, J. Ceram. Soc. Jpn. 121 (2013) 460e463. [39] W.S. Jung, Use of polymeric carbon nitride in the nitridation of gallium oxide to gallium nitride, J. Ceram. Soc. Jpn. 124 (2016) 841e843. [40] H. Zhao, M. Lei, X. Yang, Jikang Jian, X. Chen, Route to GaN and VN assisted by carbothermal reduction process, J. Am. Chem. Soc. 127 (2005) 15722e15723. [41] M.A. Reshchikov, H. Morkoç, Luminescence properties of defects in GaN, J. Appl. Phys. 97 (2005) 061301. n, M. Herrera, R. Silva, G.C. V [42] G. Guzma asquez, D. Maestre, Influence of oxygen incorporation on the defect structure of GaN microrods and nanowires. An XPS and CL study, Semicond. Sci. Technol. 31 (2016) 055006. [43] T. Kimura, S. Ootomo, T. Nomura, S. Yoshida, T. Hashizume, Solid-phase diffusion of carbon into GaN using SiNx/CNx/GaN structure, Jpn. J. Appl. Phys. 46 (2007) L224eL226. [44] L. Ramqvist, K. Hamrin, G. Johansson, A. Fahlman, C. Nordling, Charge transfer in transition metal carbides and related compounds studied by esca, J. Phys. Chem. Solids 30 (1969) 1835e1847. [45] J. Neugebauer, C.G.V.d Walle, Gallium vacancies and the yellow luminescence in GaN, Appl. Phys. Lett. 69 (1996) 503. [46] A. Sedhain, J. Li, J.Y. Lin, H.X. Jianga, Nature of deep center emissions in GaN, Appl. Phys. Lett. 96 (2010) 151902. [47] A.Y. Polyakov, M. Shin, J.A. Freitas, M. Skowronski, D.W. Greve, R.G. Wilson, On the origin of electrically active defects in AlGaN alloys grown by organometallic vapor phase epitaxy, J. Appl. Phys. 80 (1996) 6349e6354. [48] M. Julkarnain, N. Kamata, T. Fukuda, Y. Arakawa, Yellow luminescence band in undoped GaN revealed by two-wavelength excited photoluminescence, Opt. Mater. 60 (2016) 481e486.
Contents lists available at ScienceDirect
Optical Materials journal homepage: www.elsevier.com/locate/optmat
GaN nanophosphors for white-light applications Mirgender Kumar a, V.P. Singh b, Sarvesh Dubey c, Youngsuk Suh a, **, Si-Hyun Park a, * a
Department of Electronics Engineering, Yeungnam University, Gyeongsan 38541, South Korea School of Engineering, Indian Institute of Technology, Mandi 175005, India c Department of Physics, L. N. D. College, Muzaffarpur 845401, India b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 June 2017 Received in revised form 13 October 2017 Accepted 13 October 2017
GaN nanoparticles (NPs) were synthesized by carbothermal reduction combined with nitridation, using Ga2O3 powder and graphitic carbon nitride (g-C3N4) as precursors. Characterization of the NPs was performed by X-ray diffraction, scanning electron microscopy, and room-temperature photoluminescence measurements. X-ray photoelectron spectroscopy was also performed to detect the chemical states of the different species. A universal yellow luminescence (YL) band was observed from complexes of Ga vacancies with O anti-sites and of O anti-sites with C. Further increments in the C content were observed with continued growth and induced an additional blue luminescence (BL) band. Tuning of the YL and BL bands resulted in white-light emission under certain experimental conditions, thus offering a new way of employing GaN nanophosphors for solid-state white lighting. Calculations of the correlated color temperature and color-quality scale parameters confirmed the utility of the experimental process for different applications. © 2017 Elsevier B.V. All rights reserved.
Keywords: GaN nanoparticles White-light emission Correlated color temperature Color-quality scale
1. Introduction GaN shows great potential in electronics and optoelectronics because it has a wide band gap, good thermal properties, high luminescence, strong interatomic bonds, and high carrier mobility, suitable for application in light-emitting diodes (LEDs), lasers, and high-electron-mobility transistors (HEMTs) [1,2]. Ternary alloys of GaN have also reached a significant plateau in semiconductor technology and are used over a large spectral range from infrared (IR) to deep ultraviolet (UV) electromagnetic radiation [3,4]. In addition, the nontoxic nature of GaN offers an advantage over Asand phosphide-based semiconductors, which have been studied extensively in recent decades. Therefore, extensive efforts have been devoted to studying GaN in the forms of thin films and nanostructures with different morphologies, fabricated using different techniques to explore new phenomena for novel applications [5e7]. Many breakthroughs in thin-film growth have been
Abbreviations: BL, blue luminescence; CCT, correlated color temperature; CQS, color-quality scale; SEM, scanning electron microscopy; XPS, X-ray photoelectron spectroscopy; XRD, X-ray diffraction; YL, yellow luminescence. * Corresponding author.. ** Corresponding author. E-mail addresses: [email protected] (Y. Suh), [email protected] (S.-H. Park). https://doi.org/10.1016/j.optmat.2017.10.021 0925-3467/© 2017 Elsevier B.V. All rights reserved.
reported, especially for the UV/blue region; however, these remain non-cost effective for device applications. Therefore, nanoscale structures of undoped GaN may be good candidates for different applications and thus merit further research. Most GaN nanostructures have been studied, such as nanowires, nanorods, nanobelts, nanopillars, nanoparticles (NPs), and quantum dots, including special coreeshell structures [8e12]. Undoped GaN NPs have gained much attention because they have emission properties in the green, yellow, blue, and red bands of the visible spectrum, which enhance their applicability in a wide range for solid-state lighting and sensors [13e16]. These unprecedented properties have evolved from the inherent surface states accompanying vacancies (Ga vacancy, VGa, and N vacancy, VN), which act as luminescence centers and contribute to visible emission. Some surface states are generated as dangling bonds, which may act as trap centers rather than contributing as luminescence center. Some vacancy-type surface states may be occupied by residual impurities as well as anti-site atoms, which also contribute to different luminescence bands of the energy spectrum. Further, these surface states (donors or acceptors) can also form complexes with opposite-polarity states. Thus, NP surface states and complexes thereof can induce defect energy levels found between the band edges of corresponding energies and correlated well to their respective bands of luminescence [17e20]. A broad yellow luminescence (YL) band centered at 2.1e2.3 eV is a widely studied
62
M. Kumar et al. / Optical Materials 75 (2018) 61e67
visible band in the photoluminescence (PL) of GaN. An early study established that radiative transition from a shallow donor to a deep acceptor causes this YL band. In addition, by examining the broad and Gaussian nature of the emission spectrum, many studies have associated the origin of the YL band with the transition from the deep donor to shallow acceptor state, along with two transition involvement models from the shallow donor to the deep acceptor and from the deep donor to the shallow acceptor, which involve the complexes of VGa surface states and complexes of carbon defects at VN surface states with the nearest neighbors [21e24]. Some studies have also reported high-purity GaN showing a broad PL feature in the visible range, explained by de-convolution into two bands of YL centered at 2.2 eV and green luminescence (GL) centered at 2.48 eV [25,26]. The shifting of the YL band into the green region has also been observed with increases in excitation intensity, signifying that both bands correspond to the same defects with different charge or complexation states [27]. The blue luminescence (BL) band in the visible spectrum has also been studied significantly, with unstable properties assigned to specific defects [28,29]. The contribution of VGa surface states with some residual impurities has been stated as a probable reason for this band. High growth temperatures, which induce Ga desorption in N-deficient ambient atmosphere, were found to favor the generation of many VGa surface states. In addition to this correspondence, the combination of shallow donors (C occupying VGa surface states, CGa) with shallow acceptor levels (C occupying VN surface states, CN) has also been reported as a convincing cause. Some of the studies have also correlated the BL and YL band intensities with intentional or unintentional C doping, which may induce interstitial (CI) and anti-site (CN and CGa) defects [30,31]. These defects can also originate from intrinsic surface states relating to N and Ga vacancies, which can be tuned effectively in NPs. Some recent studies have reported white-light emission from GaN superlattice structures by tuning the most prominent YL and BL bands arising from interfacial state-related defects and residual impurities [32,33]. In this regard, GaN NPs might have much potential because they show excellent emission tuning capabilities by the modification of surface morphology through different synthesis processes. However, no process yet reported has successfully tuned the YL and BL bands for GaN NPs, despite their attributes as strong contenders for next-generation solid-state tunable-colortemperature white-light applications. Among all of the techniques available for synthesizing NPs, nitridation is an effective route regarding cost effectiveness, controllability of operating conditions, single-phase production, and high product purity. Generally, these reactions are performed using Ga precursors with N2 and NH3 gases at high temperatures. The low operating temperature and ease of use has made the nitridation of Ga2O3 with C and NH3 the preferred route, known as the carbothermal nitridation process [34e38]. An NH3-free nitridation process has also emerged as an alternative method with graphitic carbon nitride (gC3N4), which can be performed at a much lower temperature with better surface morphology [39,40]. Therefore, this work investigates the synthesis of GaN NPs for white-light applications by carbothermal reduction and nitridation, using Ga2O3 powder and g-C3N4 as the source materials. The synthesis was performed under different experimental conditions to tune the emission properties over a broad visible range with the variation in the concentration of residual O and C impurities. GaN NPs with tuned emission properties are good candidates for the further development of variable-color-temperature nanophosphors for white light-emitting diodes (LEDs) with good efficiency.
2. Experimental process Ga2O3 powder and g-C3N4 were used as the Ga precursor and nitridation reagent, respectively. The g-C3N4 was prepared by the thermal condensation of dicyandiamide, which was used to permit a low-temperature nitridation reaction with high-crystallinity products having good surface morphology. Ga2O3 and dicyandiamide were purchased from Sigma Aldrich with 99% purity. g-C3N4 is widely used and decomposes at 600 C into highly reactive C and N. C has strong reducing capabilities toward metal oxides, and nitrides can form in the presence of nitride ions. Hence, g-C3N4 carries the capability of reducing metal oxides into their corresponding metals followed by nitridation thereof, which solves the problem of controlling the rates of reduction and nitridation. Another advantage of using g-C3N4 is the formation of highly stable CO and N2 as byproducts. In comparison, NH3-based nitridation processes involve H2 in the reaction and degrade the morphology of the final product. A typical nitridation reaction was performed in a tubular furnace by keeping the grounded mixture of Ga2O3 and g-C3N4 (1:3 M ratio) in an alumina crucible under flowing N2 (5 mL/min) at 800 C (sample S1) or 750 C (sample S2) at atmospheric pressure. Further, samples S3 and S4 were synthesized at 750 C with the pressures of 50 Torr (6.7 kPa) and 100 Torr (13 kPa), respectively. Sample S5 was prepared with the reaction furnace pressure of 100 Torr (13 kPa) and the temperature of 650 C. Finally, samples of the GaN NPs were characterized by X-ray diffraction (XRD), scanning electron microscopy (SEM), PL spectroscopy, and X-ray photoelectron spectroscopy (XPS). 3. Results and discussion Fig. 1 shows the XRD pattern obtained using a Philips X'pert Pro
Fig. 1. XRD pattern of all samples for GaN NPs.
M. Kumar et al. / Optical Materials 75 (2018) 61e67
diffractometer with Cu Ka radiation for GaN samples using the carbothermal nitridation process with different synthesis parameters. The analysis of the XRD patterns reveals the formation of hexagonal wurtzite-structured GaN with different features corresponding to the (100), (002), (101), (102), (110), (103), (112), and (201) crystallographic planes. The structural analyses of all samples are summarized in Table 1 corresponding to the (101) dominating features. The average particle size, calculated using the Scherrer formula, for sample S2 is found to decrease because of the decrease in reaction temperature, which decreases the growth rate of NPs. However, the average particle size is increased for samples S3 and S4 because of the increase in growth rate with the decrease in ambient pressure. Similar to S2, the average particle size for S5 is also decreased compared to that of S4 with the decrease in reaction temperature. No XRD features corresponding to Ga2O3 and C3N4 are observed, confirming the full conversion of Ga2O3 to GaN. SEM image (Fig. 2(a)) using EVO-40 ZEISS depicts the surface morphology of the GaN NPs for sample S5. The reaction product comprises NPs with platelet shapes. Fig. 2(b) shows the histogram of NP size distribution for sample S5. Energy-dispersive X-ray (EDX) spectra of different samples, shown in Fig. 3, were obtained for the qualitative elemental analysis of different elements in the GaN samples synthesized under various experimental conditions. The EDX spectrum of sample S1 shows a small signal corresponding to C and O along with another intense signal for Ga and N, possibly originating from the precursors used in the reaction. Therefore, the XRD pattern and EDX elemental signals of C and O confirm their presences as residual impurities. The signal corresponding to C and O is stronger for sample S2 compared to that for S1 because of the lower reaction temperature. The same trends for C- and O-related signals are observed for samples S3 and S4, which are prepared at lower pressures than those for S1 and S2. Thus, impurities related to C and O are enhanced with decreases in reaction temperature and pressure. Quantitative elemental analyses are summarized in Table 2. Fig. 4 shows the PL spectra of the GaN NPs samples (S1, S2, S3, S4, and S5) obtained with the excitation wavelength of 325 nm from a HeeCd laser and using a photomultiplier tube as a detector. Sample S1 exhibits two different luminescence bands in the visible and UV ranges. The narrow UV luminescence band (~3.4 eV) is a typical signature of GaN, corresponding to near-band-edge emission. The broad visible band centered at 2.25 eV confirms the presence of defects within the band gap, induced by surface states (VGa and VN) and their anti-site-including complexes. This visible band is recognized as the well-known yellow band, reported by many researchers. The breadth of the band is attributed to the combination of multiple transitions, possibly including transitions from shallow donors to deep acceptors and from deep donors to shallow acceptors in this case. According to the impurities detected by EDX, O anti-site defects (O at N vacancy-type surface states, ON) and C antisite defects (C at N vacancy-type surface states, CN) probably serve as the shallow donors and acceptors, respectively, residing below the conduction band edge (CBE) at 0.30 eV and above the valence band edge (VBE) at 0.23 eV, respectively. The deep acceptor and donor may be complexes of different isolated defect levels,
Table 1 Structural parameters of the GaN NPs corresponding to (101) crystallographic plane. Sample
Position of peak 2q (deg.)
FWHM (deg.)
Average particle size (nm)
S1 S2 S3 S4 S5
36.92 36.90 36.90 36.87 36.85
0.8750 0.9487 0.7543 0.3928 0.5434
95.7 88.2 111 212.9 154.1
63
recognized as a complex of Ga vacancy-type surface states (VGa) and O anti-site defects (VGe(ON)n) (n ¼ 1e3) and a complex of C antisite defects with O anti-site defects (CNeON). VGa is an intrinsic defect in GaN and its presence is highly probable in GaN NPs; it may act as a deep acceptor at ~1.1 eV above the VBE. The VGa defects mostly exist in complexes with donors such as ON, rather than as isolated states. The conditions of VG as acceptor and ON as donor also favor the formation of (VGe(ON)n) complexes. Furthermore, the much smaller distinction between the YL and GL may also indicate the involvement of the same type of closely spaced levels, such as (VGaeON) for YL and (VGae(ON)2) for GL. The shifting of VGa closer to the VBE by forming a complex with ON may also indicate the presence of multiple types of complex, which would explain the greenish-yellow luminescence. A similar scenario also exists for the CNeON complex defect because of the different polarities of the constituent defects. However, their energy positions remain unconfirmed and require further investigation. Therefore, the VGainduced surface states and VN-induced surface states occupied by O and C are mainly responsible for the YL in the synthesized GaN NPs. The yellow emission band is stronger in the spectrum of sample S2, as well as broader and slightly shifted to a higher energy, because of the lower reaction temperature, at the expense of a weakened UV band. The formation of defects in sample S2 is clearly promoted by the decrease in temperature, which causes an increase in the intensity of the YL band. This phenomenon may occur with the decreased average size of the NPs as observed by structural analysis, which enhances the surface-to-volume ratio and creates higher numbers of surface induced impurities states. The stronger C and O signals in the EDX spectrum of sample S2 compared to those of S1 also indicate higher defect density in S2. Hence, the higher level of O in the sample may also initiate the formation of defect levels belonging to (VGae(ON)2) and (VGaeON), resulting in a shift of the luminescence peak toward a higher energy, as observed here. In response, the intensity of the UV band is quenched because of the weaker free exciton recombination with the involvement of different radiative and non-radiative defect levels, related to the enhanced surface states with the reduced surface-to-volume ratio. A new BL band appears in the spectrum of sample S3 along with a stronger YL band, as observed in the spectra of S1 and S2. The BL band intensity is lower than that of the YL band, however, and the stronger visible bands suppress the UV band. The BL band in the visible region may arise from the transition from a shallow donor defect level to a shallow acceptor defect level [41]. The existence of both BL and YL bands may also indicate the formation of a new shallow donor along with ON, expected to be near the CBE. In this case, the most probable donor state is the C anti-site defect at the Ga site (CGa) at 0.20 eV below the CBE [17]. However, this defect level requires higher formation energy. The increase in C impurities with the decrease in reaction pressure, along with the abundance of Ga vacancies, may favor the generation of CGa. Therefore, the increased level of residual impurities probably causes the strengthening of the YL band and the inception of the BL band. A further decrease in the operating pressure of the reaction to prepare sample S4 increases the intensities of both YL and BL. The BL band intensity is increased significantly compared to that of the YL band because of the larger role of CGa as a shallow donor along with the generation of interstitial O (OI) and C (CI). The presences of interstitial O and C have been confirmed experimentally in GaN, and reportedly contribute to the visible bands [42,43]. The YL band continues to increase in intensity probably because of the increasing amount of C impurities, which also induce the generation of a new mid-gap defect level corresponding to interstitial C. Further, the PL spectrum of sample S5 shows the increased intensity of both visible bands, which appear almost overlapped upon the decrease in reaction temperature, compared to those of sample S4
64
M. Kumar et al. / Optical Materials 75 (2018) 61e67
Fig. 2. SEM image (a), and the histogram of size distribution (b) for GaN NPs sample S5.
Fig. 3. EDX spectra of different NPs samples prepared with different experimental conditions.
Table 2 Percentage of different elements in GaN NP samples calculated from EDX spectra. Sample/Elements
Ga (%)
N (%)
C (%)
O (%)
S1 S2 S3 S4
81 77 72 68
14 10 8 7
2 5 9 12
3 8 11 13
with the same reaction chamber pressure. This is expected because of the increase in residual impurity levels, which enhance the generation of contributing defect levels. Moreover, the transition energy between ON and OI (~0.42 eV above the VBE) may also contribute to this broad BL, which is highly feasible with a high level of residual O [42]. The broad luminescence obtained in the white-light region is consistent with the above discussion and with previously reported results. Thus, the process of variable experimental conditions establishes GaN NPs as strong candidates for UV-excited phosphors in solid-state lighting, as they generate broad visible emission with controlled C and O contents.
Fig. 4. PL spectra with the discussed different experimental conditions.
XPS measurements were also performed for synthesized GaN NPs of sample S5 to examine the chemical state of different elements, which can confirm the existence of different defect levels. Fig. 5 shows the different XPS peaks corresponding to Ga 3d, N 1s, O 1s, and C 1s. The Ga 3d peak in Fig. 5(a) shows a slight asymmetry on the higher binding energy side. By fitting this spectrum with a Gaussian profile, it can be de-convoluted into two components, G1 (~19.7 eV) and G2 (~20.8 eV). G1 represents the dominant peak corresponding to the GaeN bond, and the higher-binding-energy peak of G2 corresponds to the GaeO bond. G2 is observed at a higher binding energy with a greater difference in the electronegativity of Ga and O compared to that of Ga and N [44]. This higherbinding-energy peak is usually attributed to oxidation states, oxynitride, GaeOH components, and hybridization states of N 2s and Ga d states. The possibility of GaeOH is negligible because H has no role in the reaction. The options of oxide and oxynitride states are also minimal, with no supporting XRD peaks. Hybridization state features also have very low possibilities in the presence of many residual O impurities. However, the present reaction conditions of low ambient pressure and temperature along with a high concentration of residual O is expected to favor O anti-site (ON) defect formation instead of the above discussed cases. This is also supported by the EDX
M. Kumar et al. / Optical Materials 75 (2018) 61e67
65
Fig. 5. XPS for synthesized GaN NPs: (a) Ga 3d peak; (b) N 1s peak; (c) O 1s peak; (d) C 1s peak.
measurements, which show a much lower concentration of N than that for Ga and suggests the partial incorporation of O into the GaN lattice by substitution of N vacancies. Finally, our reaction was performed at a lower pressure, which favors the fast carbothermal reduction of Ga2O3 and therefore removes the possibility of oxide states. If Ga2O3 were present, then it would appear in XRD because of the crystalline nature of the precursor used in the reaction. Therefore, the G2 signature is attributed to the presence of O at N sites, confirming the existence of O anti-site defects (ON) [45]. The fitted N 1s spectrum in Fig. 5(b) shows two de-convoluted peaks, N1 (~397.2 eV) and N2 (~398.2 eV). The N1 peak belongs to the NeGa bond, and the peak N2 at higher binding energy can be attributed to the CeN bond because of the higher electronegativity difference of C and N compared to that of N and Ga. Consequently, the N2 peak indicates the presence of elemental C at Ga sites, thus confirming the existence of C anti-site defects at Ga sites (CGa). Fig. 5(c) shows the O 1s spectrum of GaN NPs de-convoluted into two peaks, O1 (~530.4 eV) corresponding to the OeGa bond, and O2 (~531.1 eV) for interstitial O (OI). Hence, O1 represents the same defects state as G2 for ON, and peak O2 represents OI defects as acceptors close to the VBE [42]. After Gaussian fitting, the spectrum of C 1s shown in Fig. 5(d) is de-convoluted into three peaks, C1 (~287 eV), C2 (~286 eV), and C3 (~285 eV). The peak with lowest binding energy, C3, represents elemental C, which may originate from residual C. In addition, the diffusion of C is highly likely; C may behave as a deep acceptor in the form of an interstitial atom (CI) [43]. The high C concentration found in the EDX measurements as a residual impurity also supports the formation of interstitial states
under a low reaction chamber pressure. The higher-energy peak C2 represents the CeN bond and indicates C substitution at the Ga sites (CGa), which are also detected in the N 1s spectrum. Another higher-binding-energy peak, C1, is also shifted because of the higher difference in electronegativity between C and Ga in the CeGa bond, i.e., C substitution at N sites (CN). Fig. 6 shows the Commission Internationale de l'Eclairage (CIE) chromaticity coordinates obtained from the PL spectra of the GaN samples. The inset shows the corresponding correlated color temperature (CCT) along with the color quality scale (CQS) parameter. The coordinates of samples S1 and S2 are located in the yellow region in the chromaticity diagram, as expected. However, the coordinates of S3 are shifted to the white region because of the inception of the blue band. The coordinates of samples S4 and S5 are closer to the standard white light coordinates (0.33, 0.33) because of the reduced intensity difference between the YL and BL bands. The value of CCT for samples S1 to S5 increase continuously from 3800 to 6000 K, while the CQS value of sample S2 is lower than that of S1 and thereafter increases consistently from 89 for S3 to 95 for S5. The quality of white light of S5 appears better than those of S3 and S4 with an increased CQS, showing luminescence equivalent to that of cool daylight. Fig. 7 shows the energy level diagram of the synthesized GaN NPs, explaining the contributions of different defect levels to the different visible bands that attain broad white luminescence. As shallow donors, ON and CGa have energy levels at 0.30 and 0.20 eV below the CBE, respectively [17,46]. Indeed, the high defect concentration and lattice interaction may also induce broadening of an
66
M. Kumar et al. / Optical Materials 75 (2018) 61e67
acceptor (CN) confirms the contribution to the yellow band, which also supports the hypothesis of Glaser et al., who suggested a twostep YL transition with non-radiative capture from a shallow donor to a deep donor and radiative recombination from a deep donor to a shallow acceptor [23]. Julkarnain et al. confirmed a deep acceptor at 1 eV for the contribution to YL [48], with an earlier hypothesis of transition from a shallow donor to a deep acceptor. In our case, interstitial C is considered another deep acceptor, which evolves with high levels of residual C impurities in the samples under favorable experimental conditions. Guzm an et al. also confirmed that interstitial O at 0.42 eV above the VBE contributes to BL [42]. The intensification of both visible bands is consistent with the increasing defect states related to C and O. Our proposed band diagram also explains the YL band for GaN samples with and without C impurities. Therefore, the visible BL, GL, and YL can be converted into white-light emission. 4. Conclusions
Fig. 6. Chromaticity coordinates (CIE 1931) for the different samples, where C ¼ S1, ◄ ¼ S2, A ¼ S3, ; ¼ S4. Inset: corresponding CCT and CQS.
GaN NPs were successfully synthesized by carbothermal reduction and nitridation, as confirmed by XRD, EDX, and SEM analysis. White luminescence was also attained by varying the experimental conditions, which generated different defect levels between the band edges. The ability to tune defects related to C and O demonstrated that GaN NPs were promising candidates for nanophosphor-based white-light sources. The proposed experimental process also demonstrated the tunability of color temperature along with color quality scale for different lighting purposes. These results may enhance the further development of GaN-based devices with improved performance. Acknowledgement This work was supported by the 2014 Yeungnam University Research Grant. References
Fig. 7. Schematic of the proposed energy level diagram of the GaN NPs responsible for perceived white-light emission.
individual defect state. Therefore, the two closely spaced shallow donors can be considered localized states between 0.15 and 0.30 eV, suggesting non-radiative relaxation from the CBE and thereby supporting broad emission for different visible bands. As discussed above, the three complexes of Ga vacancies and O anti-site defects may exist in conjunction with O anti-site defects (n ¼ 1e3). However, only VGe(ON) and VGe(ON)2 can promote the emission of YL and GL, respectively. The VGa level is located 1.1 eV above the VBE, and moves to 0.9 eV after forming the complex VGe(ON) [45]. Therefore, VGe(ON)2 should shift to 0.8 eV, closer yet to the VBE. In addition, the O-containing samples exhibit higher probabilities of forming VGa and ON complexes than that for forming an isolated VGa [14]. The hypothesis of the formation of VGe(ON)2 may be supported by these facts. The possibility of a deep donor (CNeON) contributing to YL supports the scenario of two transitions creating YL, as proposed by Polyakov et al. [47]. The energy state at 0.9 eV (below the CBE) of this deep donor and the transition to a shallow
[1] T.J. Flack, B.N. Pushpakaran, S.B. Bayne, GaN technology for power electronic applications: a review, J. Electron. Mater. 45 (2016) 2673. [2] H. Morkoç, Handbook of Nitride Semiconductors and Devices, GaN-based Optical and Electronic Devices, John Wiley & Sons, 2009. [3] S. Nakamura, T. Mukai, M. Senoh, Candela-class high-brightness InGaN/AlGaN double-heterostructure blue-light-emitting diodes, Appl. Phys. Lett. 64 (1994) 1687. [4] A.G. Bhuiyan, K. Sugita, A. Hashimoto, A. Yamamoto, InGaN solar cells: present state of the art and important challenges, IEEE J. Photonics 2 (2012) 276e293. [5] E.D.L. Boulbar, C.J. Lewins, D.W.E. Allsopp, C.R. Bowen, P.A. Shields, Fabrication of high-aspect ratio GaN nanostructures for advanced photonic devices, Microelectron. Eng. 153 (2016) 132e136. [6] M.M. Rozhavskaya, W.V. Lundin, E.E. Zavarin, E.Y. Lundina, S.I. Troshkov, V.Y. Davydov, M.A. Yagovkina, P.N. Brunkov, A.F. Tsatsulnikov, Various types of GaN/InGaN nanostructures grown by MOCVD on Si(111) substrate, Phys. Status Solidi C 10 (2013) 441e444. [7] M. Nyka, W. Strek, J.M. Jabonski, L. Kepinski, R. Kudrawieca, J. Misiewicz, Synthesis, structure and optical properties of GaN nanocrystallites, Mater. Sci. Semicond. Process. 8 (2005) 511. [8] P. Corfdir, C. Hauswald, O. Marquardt, T. Flissikowski, J.K. Zettler, S.F. Garrido, L. Geelhaar, H.T. Grahn, O. Brandt, Crystal-phase quantum dots in GaN quantum wires, Phys. Rev. B 93 (2016) 115305. [9] S. Takeda, K. Ishino, Y. Inoue, A. Ishida, H. Fujiyasu, H. Kominami, H. Mimura, Y. Nakanishi, S. Sakakibara, Growth of pillarlike GaN nanostructures, Jpn. J. Appl. Phys. 44 (2005) 5664. [10] Z. Zang, X. Zeng, J. Du, M. Wang, X. Tang, Femtosecond laser direct writing of microholes on roughened ZnO for output power enhancement of InGaN lightemitting diodes, Opt. Lett. 41 (2016) 3463. [11] L. Zhang, C. Liu, White light emission from GaN stack layers doped by different rare-earth metals, Jpn. J. Appl. Phys. 54 (2015) 02BA03. [12] M. Nyk, R. Kudrawiec, W. Strek, J. Misiewicz, Synthesis and optical properties of Eu3þ and Tb3þ doped GaN nanocrystallite powders, Opt. Mater. 28 (2006) 767. [13] Y.H. Ko, J. Song, B. Leung, J. Han, Y.H. Cho, Multi-color broadband visible light source via GaN hexagonal annular structure, Sci. Rep. 4 (2014) 5514.
M. Kumar et al. / Optical Materials 75 (2018) 61e67 [14] X. Li, P.W. Bohn, J.J. Coleman, Impurity states are the origin of yellow-band emission in GaN structures produced by epitaxial lateral overgrowth, Appl. Phys. Lett. 75 (1999) 4049. [15] C. Li, Z. Zang, W. Chen, Z. Hu, X. Tang, W. Hu, K. Sun, X. Liu, W. Chen, Highly pure green light emission of perovskite CsPbBr3 quantum dots and their application for green light-emitting diodes, Opt. Express 24 (2016) 15071. [16] X. Shi, M. Peng, J. Kou, C. Liu, R. Wang, Y. Liu, J Zhai, A Flexible GaN nanowire array-based schottky-type visible light sensor with strain-enhanced photoresponsivity, Adv. Electron. Mater. 1 (2015) 1500169. [17] A.F. Wright, Substitutional and interstitial carbon in wurtzite GaN, J. Appl. Phys. 92 (2012) 2575. [18] A. Sedhain, J. Li, J.Y. Lin, H.X. Jianga, Nature of deep center emissions in GaN, Appl. Phys. Lett. 96 (2010) 151902. [19] N.T. Son, C.G. Hemmingsson, T. Paskova, K.R. Evans, A. Usui, N. Morishita, n, Identification of the gallium T. Ohshima, J. Isoya, B. Monemar, E. Janze vacancyeoxygen pair defect in GaN, Phys. Rev. B 80 (2009) 153202. [20] B. Schwenzer, S. Keller, L. Loeffler, R. Seshadri, F.F. Lange, S.P. Baars, U.K. Mishra, Group-III nitride nanoparticles-synthesis and photoluminescence studies, in: Conference Proceeding of IEEE-NANO, 2003, pp. 12e14. [21] T. Ogino, M. Aoki, Jpn. Mechanism of yellow luminescence in GaN, J. Appl. Phys. 19 (1980) 2395. [22] D.M. Hofmann, D. Kovalev, G. Steude, B.K. Meyer, A. Hoffmann, L. Eckey, R. Heitz, T. Detchprom, H. Amano, I. Akasaki, Properties of the yellow luminescence in undoped GaN epitaxial layers, Phys. Rev. B 52 (1995) 16702e16706. [23] E.R. Glaser, T.A. Kennedy, K. Doverspike, L.B. Rowland, D.K. Gaskill, J.A. Freitas, M.A. Khan, D.T. Olson, J.N. Kuznia, D.K. Wickenden, Optically detected magnetic resonance of GaN films grown by organometallic chemical-vapor deposition, Phys. Rev. B 51 (1995) 13326e13336. [24] E. Calleja, F.J. Sanchez, D. Basak, M.A.S. Garcia, E. Munoz, I. Izpura, F. Calle, J.M.G. Tijero, J.L.S. Rojas, B. Beaumont, P. Lorenzini, P. Gibart, Yellow luminescence and related deep states in undoped GaN, Phys. Rev. B 55 (1997) 4689. [25] M.A. Reshchikov, H. Morkoc, S.S. Park, K.Y. Lee, Yellow and green luminescence in a freestanding GaN template, Appl. Phys. Lett. 78 (2001) 3041e3043. [26] M.A. Reshchikov, D.O. Demchenko, A. Usikov, H. Helava, Y. Makarov, Carbon defects as sources of the green and yellow luminescence bands in undoped GaN, Phys. Rev. B 90 (2014) 235203e235211. [27] M.A. Reshchikov, H. Morkoc, S.S. Park, K.Y. Lee, Two charge states of dominant acceptor in unintentionally doped GaN: evidence from photoluminescence study, Appl. Phys. Lett. 81 (2002) 4970e4972. [28] M. Gopalakrishnan, V. Purushothaman, V. Ramakrishnan, G.M. Bhaleraoc, K. Jeganathan, The effect of nitridation temperature on the structural, optical and electrical properties of GaN nanoparticles, CrystEngComm 16 (2014) 3584e3591. [29] Y. Yang, V.J. Leppert, S.H. Risbud, B. Twamley, P.P. Power, H.W.H. Lee, Blue luminescence from amorphous GaN nanoparticles synthesized in situ in a polymer, Appl. Phys. Lett. 74 (1999) 2262e2264. ndez, R.M. Pe rez, B.M. Monroy, A.E. Esquivel, [30] G. Santana, O. Melo, J.A. Herna pez, F. Moure, L.A. Hern M.L. Lo andez, G.C. Puente, Photoluminescence study of gallium nitride thin films obtained by infrared close space vapor transport,
67
Materials 6 (2013) 1050e1060. [31] J. Kang, Y. Shen, Z. Wang, Effects of residual C and O impurities on photoluminescence in undoped GaN epilayers, Mater. Sci. Eng. B91e92 (2002) 303e307. [32] N.B. Sedrine, T.C. Esteves, J. Rodrigues, L. Rino, M.R. Correia1, M.C. Sequeira, A.J. Neves, E. Alves, M. Bockowski, P.R. Edwards, K.P. O'Donnell, K. Lorenz, T. Monteiro, Photoluminescence studies of a perceived white light emission from a monolithic InGaN/GaN quantum well structure, Sci. Rep. 5 (2015) 13739e13741. [33] M.L. Lee, Y.H. Yeh, S.J. Tu, P.C. Chen, W.C. Lai, J.K. Sheu, White emission from non-planar InGaN/GaN MQW LEDs grown on GaN template with truncated hexagonal pyramids, Opt. Express 23 (2015) A401eA412. [34] C.M. Balkas¸, R.F. Davis, Synthesis routes and characterization of high-purity, single-phase gallium nitride powders, J. Am. Ceram. Soc. 79 (1996) 2309e2312. [35] T.H. Kim, S. Choi, D.W. Park, Thermal plasma synthesis of crystalline gallium nitride nanopowder from gallium nitrate hydrate and melamine, Nanomaterials 6 (2016) 38. [36] V. Ganesh, S. Suresh, M. Balaji, K. Baskar, Synthesis and characterization of nanocrystalline gallium nitride by nitridation of Ga-EDTA complex, J. Alloys Compd. 498 (2010) 52e56. [37] W.S. Jung, Reaction intermediate(s) in the conversion of h-gallium oxide to gallium nitride under a flow of ammonia, Mater. Lett. 57 (2002) 110e114. [38] W.S. Jung, Reaction mechanism for the ammonolysis of b-gallium oxide to gallium nitride, J. Ceram. Soc. Jpn. 121 (2013) 460e463. [39] W.S. Jung, Use of polymeric carbon nitride in the nitridation of gallium oxide to gallium nitride, J. Ceram. Soc. Jpn. 124 (2016) 841e843. [40] H. Zhao, M. Lei, X. Yang, Jikang Jian, X. Chen, Route to GaN and VN assisted by carbothermal reduction process, J. Am. Chem. Soc. 127 (2005) 15722e15723. [41] M.A. Reshchikov, H. Morkoç, Luminescence properties of defects in GaN, J. Appl. Phys. 97 (2005) 061301. n, M. Herrera, R. Silva, G.C. V [42] G. Guzma asquez, D. Maestre, Influence of oxygen incorporation on the defect structure of GaN microrods and nanowires. An XPS and CL study, Semicond. Sci. Technol. 31 (2016) 055006. [43] T. Kimura, S. Ootomo, T. Nomura, S. Yoshida, T. Hashizume, Solid-phase diffusion of carbon into GaN using SiNx/CNx/GaN structure, Jpn. J. Appl. Phys. 46 (2007) L224eL226. [44] L. Ramqvist, K. Hamrin, G. Johansson, A. Fahlman, C. Nordling, Charge transfer in transition metal carbides and related compounds studied by esca, J. Phys. Chem. Solids 30 (1969) 1835e1847. [45] J. Neugebauer, C.G.V.d Walle, Gallium vacancies and the yellow luminescence in GaN, Appl. Phys. Lett. 69 (1996) 503. [46] A. Sedhain, J. Li, J.Y. Lin, H.X. Jianga, Nature of deep center emissions in GaN, Appl. Phys. Lett. 96 (2010) 151902. [47] A.Y. Polyakov, M. Shin, J.A. Freitas, M. Skowronski, D.W. Greve, R.G. Wilson, On the origin of electrically active defects in AlGaN alloys grown by organometallic vapor phase epitaxy, J. Appl. Phys. 80 (1996) 6349e6354. [48] M. Julkarnain, N. Kamata, T. Fukuda, Y. Arakawa, Yellow luminescence band in undoped GaN revealed by two-wavelength excited photoluminescence, Opt. Mater. 60 (2016) 481e486.