CRYSTAL GROWTH & DESIGN
Synthesis of GaN Nanocrystals through Phase Transition from Hexagonal to Cubic Structures upon Laser Ablation in Liquid P. Liu, Y. L. Cao, H. Cui, X. Y. Chen, and G. W. Yang* State Key Laboratory of Optoelectronic Materials and Technologies, Institute of Optoelectronic and Functional Composite Materials, School of Physics Science & Engineering, Zhongshan UniVersity, Guangzhou 510275, P. R. China
2008 VOL. 8, NO. 2 559–563
ReceiVed June 29, 2007; ReVised Manuscript ReceiVed October 16, 2007
ABSTRACT: The phase transition from hexagonal to cubic GaN has been observed upon pulsed-laser ablation of hexagonal GaN powders in a liquid at room temperature and ambient pressure. At the same time, GaN nanocrystals are synthesized through this hexagonal-to-cubic phase transition. Cathodoluminescence spectroscopy is employed to characterize the luminescence of the synthesized GaN nanocrystals. First-principle calculations are used to clarify the physical and chemical mechanisms of the phase transition from hexagonal to cubic GaN upon pulsed-laser induced solid–liquid interface reaction.
I. Introduction Group III metal nitrides nanomaterials and nanostructures, such as particles, wires, and tubes, have become the focus of intensive research because of their unique applications in mesoscopic physics and in the fabrication of nanodevices.1 GaN is a wide bandgap compound semiconductor and a greatly promising candidate for developing high efficiency blue laser devices at room temperature. Thus, GaN nanocrystals have received considerable interest because of the great potential applications in micro- and nano-optoelectronics.2–4 It is well known that there are two available structures of GaN: one is the thermodynamically stable phase of the hexagonal structure and the other one is the metastable phase of the cubic structure. In other words, the ground-state structure of GaN is the hexagonal wurtzite (WZ) phase under ambient conditions, and the formation of the metastable cubic zinc-blende (ZB) phase of GaN is difficult to obtain under thermodynamic equilibrium conditions. However, the cubic GaN (c-GaN) has advantageous physical properties compared to that of the hexagonal GaN (hGaN), such as easier p-type doping and easier cleaving for laser facets.4 Therefore, numerous theoretical and experimental studies of the hexagonal-to-cubic phase transformation of GaN have been reported, their aim being the preparation of c-GaN.5,6 For instance, a high pressure can induce the hexagonal phase of GaN to transform into the cubic phase.7,8 Therefore, the hexagonal-to-cubic phase transformation of GaN seems an effective route to prepare the metastable cubic ZB phase from the hexagonal WZ phase.6 Recently, we have developed a laserbased technique to synthesize nanomaterials and nanostructures with metastable phases,9,10 that is, pulsed-laser induced liquid–solid interface reaction (PLIIR). Using PLIIR, we have synthesized a series of nanocrystals of diamond and related materials. In this contribution, we report that the phase transition from hexagonal to cubic GaN is achieved upon PLIIR at room temperature and ambient pressure. At the same time, GaN nanocrystals are synthesized through this hexagonal-to-cubic phase transition.
II. Experimental Section The details of the experimental procedures have been described in a previous work.10 In this work, h-GaN powders with a purity of 98% * Corresponding author. E-mail:
[email protected].
Figure 1. XRD patterns of (a) the starting h-GaN precursors and (b) the synthesized products. are first rubbed in an agate kettle. Then, the rubbed powders are mixed with ethanol (purity > 99.7%) to be the reactive liquid in PLIIR. The second harmonic is produced by a Q-switch Nd:YAG laser device with a wavelength of 532 nm, a pulse width of 10 ns, a repetition frequency of 3 Hz, and a power density of 1010 W/cm2. After several hours of laser irradiation, the reactive liquid is transferred to a vacuum oven for drying. Finally, the synthesized products are collected as the sample in our case. X-ray diffraction (XRD) is performed with a Rigaku D/Max-IIIA X-ray diffractometer with Cu KR radiation (λ ) 1.54056 Å) at a scanning rate of 0.07° s-1, and transmission electron microscopy (TEM) is carried out with a JEOL JEM-2010H instrument, equipped with an energy dispersive X-ray spectrometer (EDS). These techniques are used to identify the structure and morphology of the products. Cathodoluminescence (CL) spectroscopy, which is carried out with a Gatan MonoCL3 system attached to a field emission scanning electron microscope, is employed to characterize the luminescence of the synthesized GaN nanocrystals.
III. Results and Discussion A. XRD Analysis. Figure 1 shows the XRD patterns of the h-GaN precursors and the synthesized products. Clearly, we can see that the main diffraction peaks that appear in the h-GaN precursors belong to the h-GaN phase with diffraction data of a WZ structure (JCPDS card File No. 500792),11 and the O-Ga peak belongs to the impurity metal Ga. However, in the XRD pattern of the synthesized products, three new peaks appear at
10.1021/cg0705963 CCC: $40.75 2008 American Chemical Society Published on Web 01/19/2008
560 Crystal Growth & Design, Vol. 8, No. 2, 2008
34.9°, 40.44°, and 58.53°, which are identified specifically to be the (111), (200), and (220) crystalline planes of the c-GaN phase with a ZB structure (JCPDS card File No. 882364).12 Additionally, the peak at 32.86° originates from the silicon support, and the other peaks are from h-GaN powders and metal Ga with an orthorhombic structure in the XRD pattern of the synthesized products. Therefore, the XRD data show that the c-GaN phase is synthesized upon PLIIR. Note that, in our laboratory, the pulsed laser is usually focused in a very small area. Thus, it seems difficult to irradiate all the h-GaN precursors with the pulsed laser during synthesis. Therefore, we have not achieved a high efficiency of the hexagonal-to-cubic phase transition by using the present technique based on PLIIR. In other words, it is currently difficult to synthesize the pure c-GaN nanocrystals by the hexagonalto-cubic phase transition upon PLIIR. As we see in Figure 1, the synthesized products contain c-GaN and h-GaN. Therefore, the high phase-transformation efficiency could be obtained by irradiating the h-GaN precursors more exhaustively or by modifying the h-GaN precursors to be a solid target. B. TEM Analysis. To gain further information on the synthesized products, a careful TEM analysis was carried out. Figure 2 shows a bright-field TEM image of the synthesized products, the corresponding selected area electron diffraction pattern (SAED), and a high-resolution TEM image. In Figure 2a, we can see that the resultant products are quasi-spherical nanoparticles. Also, some amorphous materials are dispersed in the synthesized products. With the EDS analysis, we find that the amorphous phase is a mixture of amorphous carbon (which could come from the pure ethanol) and a small amount of GaN. However, these amorphous materials have no influence on our further analysis. The quasi-spherical nanoparticles’ sizes are measured to be in the range of 50-150 nm. A single quasispherical nanoparticle is shown in Figure 2b, and the corresponding SAED pattern is shown in the inset of Figure 2b. On the basis of a careful analysis, the SAED pattern is indexed exactly to be [111], [220], [131], and [311] of the c-GaN phase with a ZB structure. Furthermore, the corresponding highresolution TEM image, shown in Figure 2c, shows the lattice planes with a spacing of 0.258 nm, which are assigned to be the (111) planes of c-GaN. The twining planes and some staking faults are visible in the high-resolution TEM image (they are marked by a circle in Figure 2c), and the corresponding SAED analysis, shown in Figure 2d, indicates that there are two types of lattices which can be indexed by two twining planners: one type is a hexagonal lattice with a WZ structure, and the other type is the (111) planes of c-GaN. Therefore, these analyses show that the polytype modulated (WZ and ZB) twined crystals are present in the synthesized products, and the phase transformation from h-GaN to c-GaN seems incomplete in some nanocrystals. Accordingly, these results clearly show that the c-GaN nanocrystals are synthesized through the pulsed-laser ablation of h-GaN powders in liquid. On the other hand, we have analyzed the particles with edges and corners visible in Figure 2a. Apparently, they are usually large (more than 200 nm) and have a WZ structure. Figure 2e shows a TEM image and the corresponding SAED pattern (shown in the inset) of one such particle with edges and corners, and the SAED can be indexed exactly to the [101], [102], and [103] of a WZ structure. Moreover, the corresponding highresolution TEM image, shown in Figure 2f, shows the lattice planes with a spacing of 0.244 nm, which are assigned to be the (101) planes of h-GaN. Therefore, based on these experi-
Liu et al.
mental analyses, we can conclude that most of the c-GaN nanoparticles in the synthesized products have a quasi-spherical shape and are small, and the h-GaN particles that have not undergone the phase transformation in the synthesized products keep the shape with edges and corners and are large. Thus, it is easy to distinguish the c-GaN nanoparticles from the h-GaN particles in the synthesized products. C. CL Spectrum. Figure 3a shows an SEM image of several c-GaN nanocrystals gobbets prepared on a silicon substrate, and the corresponding CL image and CL spectrum are shown in panels (b) and (c) of Figure 3, respectively. A strong CL peak at 376 nm at room temperature is displayed in Figure 3c, which suggests that the energy Eg of the CL peak is 3.29 eV. Thus, this result is in agreement with the previous study of c-GaN (3.25 eV).6,13 Note that the energy Eg of the high-energy CL band of h-GaN is 3.45 eV.3,13 Therefore, the CL spectrum obtained from the pure h-GaN crystals shows the CL peak at 362 nm (Eg ) 3.47 eV) at room temperature, which is in agreement with the literature aforementioned. Accordingly, these results definitely show that the CL peak at 376 nm is from the c-GaN nanocrystals. D. Synthesis Mechanisms. PLIIR is developed from laser ablation in liquid to synthesize compounds in a metastable phase with a size on the scale of nanometers.10 Generally, PLIIR is a very fast and far from equilibrium process; thus, most of the stable and metastable phases forming at the initial, intermediate, and final stages of the conversion may be present in the final products, especially the metastable intermediate phase.14 At the initial stage of the laser ablation, the species ejected from the solid surface would form a dense region, that is, a plasma plume, in the vicinity of the solid–liquid interface because of the confinement effect of the liquid.14 Because the plasma plume is strongly confined in the liquid, and the compressibility of the liquid is considered several orders of magnitude higher than that of the corresponding gas, the liquid structure breaks the plasma expansion to form an adiabatic region.15 In this adiabatic region, a shock wave is created with a supersonic velocity, and it induces an extra pressure, called laser-induced pressure, in the plasma plume. Furthermore, the laser-induced pressure induces a temperature increase in the plasma plume. Therefore, compared with the plasma plume from a laser ablation in gas or vacuum, the plasma plume formed in PLIIR is at a higher temperature, higher pressure, and higher density state. As a result of the confinement effect of the liquid, the quenching time of the plasma plume becomes so short that the metastable phase forming at the intermediate stage of the process can be frozen in the synthesized final products. Accordingly, we can divide PLIIR into three stages: the formation, transformation, and condensation of the plasma plume from the laser ablation of solids in liquids.10 In our case, the hexagonal-to-cubic phase transition is achieved via the three stages above. At the stage of the plasma plume formation, the laser-induced plasma, which includes a lot of h-GaN species with a WZ structure, is first generated at the liquid–solid interface when the pulsed laser ablates the h-GaN powders in the liquid. Then, the laser-induced plasma is driven into a high temperature, high pressure, and high density state because of the laser irradiation and laser-induced pressure mentioned above, which drives the plasma plume into the stage of the transformation process. Because the c-GaN phase is thermodynamically stable under conditions of high temperature and pressure, these h-GaN species with a WZ structure in the plasma plume can transform into the c-GaN phase with a ZB structure. In detail, the (0002) basic
GaN Nanocrystals Synthesis through Phase Transition
Crystal Growth & Design, Vol. 8, No. 2, 2008 561
Figure 2. (a) Bright-field TEM image of the synthesized products. (b) Bright-field TEM image of an individual nanocrystal. The corresponding SAED is shown in the inset. (c) Corresponding high-resolution TEM image of the single nanocrystal shown in panel b. (d) Corresponding SAED pattern of the twining planes in the nanocrystal shown in panelc. The circles indicate the c-GaN phase. (e) Bright-field TEM image of a single h-GaN particle with edges and sides. The corresponding SAED is shown in the inset. (f) Corresponding high-resolution TEM image of one corner of the single particle shown in panel e.
plane in the h-GaN structure has a compression displacement without diffusion, and the neighbor atoms on the (0002) basic plane adjust their displacement by an average of 0.7% during the phase transformation. Finally, the (0002) basic plane in the h-GaN structure transforms into the (111) plane of the
ZB structure.16 Figure 4 shows the mechanism of the phase transition. At the stage of the plasma plume condensation, the c-GaN phase with a ZB structure from the h-GaN species nucleates and grows during the rapid quenching of the plasma plume in the confined liquid. Because the growth time
562 Crystal Growth & Design, Vol. 8, No. 2, 2008
Liu et al.
Figure 3. (a) SEM image of several c-GaN nanocrystals. (b) Corresponding CL image. (c) Corresponding CL spectrum. (d) CL spectrum obtained from the pure h-GaN crystals.
Figure 4. Schematic illustration of the mechanism of the h-GaN-to-c-GaN phase transition upon PLIIR.
(plasma quenching time) of the nuclei is very short, the diameter of the grown crystals is usually at the nanometer scale.15 Figure 5 shows the comparison of h-GaN (solid line) and c-GaN (dashed line) total energies, which suggests that the difference between the equilibrium total energies of the two phases is very small. This result implies that the hexagonal-tocubic phase transition of GaN is possible upon a far from equilibrium process such as PLIIR.17,18 Accordingly, the formation of c-GaN upon PLIIR seems to originate from the hexagonal-to-cubic phase transition of GaN. Therefore, PLIIR
obviously provides an efficient route to synthesize c-GaN nanocrystals.
IV. Conclusion In summary, we have performed the hexagonal-to-cubic phase transition of GaN upon pulsed-laser ablation in a liquid. Then, we have synthesized GaN nanocrystals via the hexagonal-tocubic phase transition. Interestingly, a strong CL peak at 376 nm at room temperature is observed in the CL spectrum of the synthesized c-GaN nanocrystals. Additionally, the synthesis
GaN Nanocrystals Synthesis through Phase Transition
Crystal Growth & Design, Vol. 8, No. 2, 2008 563
References
Figure 5. Total energy of c-GaN and h-GaN calculated by first-principle calculations, in which an accurate full-potential, linearized, and augmented plane wave method is used, and the exchange and correlation effects are treated within the generalized gradient approximation.
mechanisms are proposed to address the formation of c-GaN nanocrystals upon PLIIR. Acknowledgment. The National Natural Science Foundation of China (50525206, 10474140, and U0734004) and the Ministry of Education (106126) supported this work.
(1) Duan, X. F.; Huang, Y.; Cui, Y.; Wang, J. F.; Lieber, C. M. Nature 2001, 409, 66. (2) As, D. J.; Richter, A.; Busch, J.; Lubbers, M.; Mimkes, J.; Lischka, K. Appl. Phys. Lett. 2000, 76, 13. (3) Li, S.; Ouyang, C. Y. Phys. Lett. A 2005, 336, 145. (4) Wang, W. Y.; Xu, Y. P.; Zhang, D. F.; Chen, X. L. Mater. Res. Bull. 2001, 36, 2155. (5) Blumenau, A. T.; Fall, C. J.; Elsner, J.; Jones, R.; Heggie, M. I.; Frauenheim, T. Phys. Status Solidi C 2003, 0, 1684. (6) Dhara, S.; Datta, A.; Wu, C. T.; Lan, Z. H.; Chen, K. H.; Wang, Y. L.; Hsu, C. W.; Shen, C. H.; Chen, L. C.; Chen, C. C. Appl. Phys. Lett. 2004, 84, 5473. (7) Purdy, A. P. Chem. Mater. 1999, 11, 1648. (8) Al-Sharif, A. I. Solid State Commun. 2005, 135, 515. (9) Yang, G. W.; Wang, J. B.; Liu, Q. X. J. Phys.: Condens. Matter 1998, 10, 7923. (10) Yang, G. W. Prog. Mater. Sci. 2007, 52, 648. (11) Balkas, C.; Basceri, C.; Divas, R. Powder Diffr. 1995, 10, 226. (12) Wright, A. F.; Nelson, J. S. Phys. ReV. B 1995, 51, 7866. (13) Flores, G. R.; Contreras, H. N.; Matinez, A. L.; Powell, R. C.; Greene, J. E. Phys. ReV. B 1994, 50, 8433. (14) Berthe, L.; Fabbro, R.; Peyer, P.; Bartnicki, E. J. Appl. Phys. 1999, 85, 7552. (15) Wang, C. X.; Liu, P.; Cui, H.; Yang, G. W. Appl. Phys. Lett. 2005, 87, 201913. (16) Yang, G. W.; Wang, J. B. Appl. Phys. A. 2001, 72, 475. (17) Wang, C. X.; Yang, Y. H.; Liu, Q. X.; Yang, G. W.; Mao, Y. L.; Yan, X. H. Appl. Phys. Lett. 2003, 84, 1471. (18) Liu, Q. X.; Yang, G. W.; Zhang, J. X. Chem. Phys. Lett. 2003, 373, 57.
CG0705963
