Wide band-gap materials such as diamond or III-nitride semiconductors (GaN, AlN and their alloys) are attractive for applications such as ultraviolet light-emitting diodes, lasers and high-power photonic devices1,2,3,4,5. For shorter wavelength optoelectronic applications, rare-earth doped III-nitrides also have significant potential for high-efficiency luminescent applications including photoemitters, phosphors and scintillators over wide wavelength ranges6,7. Wurtzite-type AlN (w-AlN) has one of the widest direct band-gaps of the nitrides (6.2 eV), a high thermal conductivity (~300 W m−1 K−1) and excellent physical and chemical stability8,9 and could thereby enable new electroluminescent properties in the ultraviolet or visible-light wavelength ranges. However, the key issues that must be overcome for w-AlN to realize its full potential are precise impurity control and effective, stable doping of luminous elements into the host lattices. Up to now, sublimation growth10 or hydride vapor phase epitaxy methods11 have been used to produce w-AlN single crystals. Although these methods are practically useful to synthesize millimeter size single crystals, it is difficult to control the carbon and oxygen impurity levels. Moreover, it is difficult to dope rare-earth elements into the host lattice due to the large size mismatch. As an alternative, therefore, ion implantation has been extensively used for the rare-earth doping of thin films12,13,14,15. However the resultant defect structure involving the dopants is uncontrollable and meta-stable defect structures may be introduced that are detrimental to luminescent efficiency.

Here we report the stable doping of single isolated Ce atoms into w-AlN single crystals. To overcome the above issues, we use a method involving a reactive flux, a temperature gradient at high pressure and temperature16,17,18, which results in high-quality w-AlN:Ce single crystals with visible-light luminescence (see Figure 1(a)). To directly identify the atomic sites of individual single Ce dopants within w-AlN single crystals, we use atomic-resolution scanning transmission electron microscopy (STEM) in the annular dark-field (ADF) imaging mode, which has a strong dependence on atomic number (Z-contrast)19 and is a powerful method to visualize single heavy dopants18. Combining with systematic first-principles calculations, we uncover the atomic structure of high-efficiency Ce centers in the wide band-gap w-AlN.

Figure 1
figure 1

Photoluminescence and XANES spectra obtained from w-AlN:Ce single crystals.

(a), (b) Pink-colored luminous image obtained by mercury lamp excitation and the PL spectrum obtained from w-AlN:Ce single crystals, where He-Cd laser (325 nm) was used for the excitation source of the PL measurement. (c) Ce M4,5 XANES spectra obtained from the w-AlN:Ce single crystals and reference samples of CeF3 (Ce3+) and CeO2 (Ce4+).

The present w-AlN:Ce single crystals were synthesized by the temperature gradient method at 6.5 GPa and 1600°C. The resulting sub-millimeter w-AlN:Ce single crystals (~0.2 mm cube) show optical transparency and the band edge luminescent feature of w-AlN is observed at 210 nm in cathodoluminescence (CL) spectroscopy, suggesting that the impurity level is sufficiently suppressed. Figure 1(a) shows a photoluminescence (PL) image of w-AlN:Ce single crystals. The homogeneous pink-colored luminescence shows that the Ce dopants are uniformly doped throughout the single crystals. A single broadband peak, observed at 600 nm in the PL spectrum of Figure 1 (b), is assigned to the 4f – 5d electron transition of Ce3+ dopants in a bulk material. The valence state of Ce dopants in the w-AlN single crystals was evaluated by Ce M4,5 X-ray absorption near edge structure (XANES). Two specific peaks in Figure 1(c) are assigned to the Ce M5 (879 eV) and M4 (897 eV) edges. By comparing to the standard XANES spectra of Ce3+ in CeF3 and Ce4+ in CeO2, the valence state of the Ce dopants in w-AlN is determined to be 3+.

Figure 2(a) shows a typical atomic-resolution ADF STEM image of w-AlN:Ce viewed along the direction. The scattering strength of nitrogen is relatively weak, but the atomic dumbbell consisting of Al and N atoms is clearly resolved (the projected spacing is 1.1 Å), as is evident in Figure 2(b). Thus, we can directly determine the atomic site of the Ce dopants from the images, whether they are on Al, N or interstitial sites. As marked by the arrows in Figure 2(a), the atomic columns containing Ce dopants are observed as slightly brighter spots and one can see that the Ce dopants are dispersed as single atoms. This isolation of single Ce dopants is critical for the optical luminescence because dopant-clusters lose radiative recombination efficiency, in particular for large clusters, owing to their local metallic properties. The present rare-earth doping process in w-AlN is thus highly effective and all Ce dopants should contribute equally to the pink-colored luminescence. Note that, even after the extensive observations, we could not observe any rare-earth dimer- or trimer-clusters, as previously proposed12,14, suggesting that these clusters may be meta-stable defect structures. In the high-magnification atomic-resolution ADF STEM image of Figure 2(b), the single Ce dopant evidently substitutes for the Al site without apparent atomic displacements. This bright feature at the Al site is quantitatively reproduced in the simulated image (Figure 2(c)), where a single Ce atom has been substituted for an Al atomic site.

Figure 2
figure 2

Experimental and simulated Atomic-resolution Z-contrast STEM images.

(a) Atomic-resolution ADF STEM image of w-AlN:Ce viewed along the direction. The brighter Al columns marked by arrows contain single Ce atoms. (b) High magnification ADF STEM image, where 10 sequentially acquired frames are averaged after image alignment by cross-correlation; no other image processing is applied. (c) Simulated image based on the neutral CeAl defect structure with an inserted schematic showing Al and N positions. The specimen thickness is estimated to be 5.3 ± 0.6 nm by quantification of the ADF image25 and the Ce atom position is assumed to be at a depth of 1.6 nm. The scale bars are (a) 1 nm and (b) 0.2 nm.

In order to theoretically estimate the stable Ce atom configurations in the w-AlN lattice, we performed systematic first-principles calculations using the Heyd-Scuseria-Ernzerhof (HSE06) hybrid functional20,21, which has been applied to Ce dopants in c-BN18 as well as a variety of solids including rare-earth compounds22,23,24. The formation energy (Ef) of a defect is written as

where denotes the total energy of the supercell (192 atoms) with a defect in charge state q, ni is the number of the constituent atoms of i-type and μi and are the atomic chemical potential and the Fermi level. The chemical potentials for the w-AlN:Ce single crystals are not clear under the present synthesis conditions and therefore we treated the chemical potentials as variables. μAl, μCe and μN are considered to range between the following extreme conditions: the Al-rich limit (μAl = μAl(bulk), μN = 1/2μN2 + ΔEf (AlN) and μCe = μCe(bulk)) and the N-rich limit (μN = 1/2μN2, μAl = μAl(bulk) + ΔEf (AlN) and μCe = μCe(bulk) + ΔEf (CeN)), where μAl(bulk), μCe(bulk) and μN2 are the chemical potentials of the Al and Ce crystals and a N2 molecule given by the calculated total energies and ΔEf (AlN) and ΔEf (CeN) are the calculated formation energies of w-AlN and CeN, −3.14 and −2.97 eV, respectively. In this calculation, we consider several defect structures and complexes: (1) Ce substitutions on the Al and N sites, CeAl and CeN; (2) Ce as an interstitial at the octahedral site, Cei; (3) Ce coupled with an Al-vacancy, CeAl-ViAl (i = a, b, or c, where CeAl and VAl are located on the same c-plane or VAl is located on the c-plane below or above CeAl; and (4) Ce coupled with a N-vacancy, CeAl-VjN ( or , where CeAl and VN are located along the direction parallel to or nearly perpendicular to the c axis). In addition, isolated Al and N vacancies (VAl and VN) were considered as the dominant native defects in w-AlN.

Figures 3(a) and (b) show the formation energies of several defect structures and complexes considered here as a function of the Fermi level. The formation energies of VN and VAl (dashed lines in Figures 3(a) and (b)) are negative in the lower Fermi levels (< 1.2 eV) at the Al-rich limit and in the higher Fermi levels (> 4.4 eV) at the N-rich limit, respectively. The Fermi level cannot be located in these ranges because of carrier compensation associated with the spontaneous formation of negatively charged VAl and positively charged VN. The XANES measurement suggests the presence of Ce3+ in the w-AlN:Ce single crystals. From the formation energy and one-electron structure analyses, the energetically stable defect structures that involve Ce3+ are (1) neutral CeAl, or (CeAl-ViAl)3− at the Al-rich limit and (2) neutral CeAl or (CeAl-ViAl)3− at the N-rich limit, depending on the Fermi level. The (CeAl-ViAl)3− defect complexes are the most energetically preferable structures at a high Fermi level, but after structural relaxation the Ce atoms show large atomic shifts away from the Al sites (0.61 Å toward VAl, where i = c is the most stable defect structure in (CeAl-ViAl)), which is inconsistent with the ADF STEM observations. The defect complex also shows a lower formation energy than CeAl, but the condition is limited to very high Fermi levels (> 5.2 eV) and only found at the Al-rich limit. In the relaxed defect structure, the Ce atom exhibits a substantial atomic shift toward the N vacancy site from the Al site (0.23 Å). On the other hand, neutral CeAl is the most energetically stable defect structure under the remaining wide range of conditions (the chemical potential only slightly affects the formation energy). As seen in the relaxed structure (Figures 3(c) and (d)), the Ce atom is located at the Al site without apparent atomic displacement, which is in good agreement with the observed ADF STEM images. Thus, we conclude that the atomic structure luminescent center in w-AlN is neutral CeAl.

Figure 3
figure 3

Theoretical formation energies of Ce point-defect structures in w-AlN.

Formation energies of CeAlq, CeNq, Ceiq, (CeAl-ViAl)q and (CeAl-VjN)q defect structures along with VAl and VN as a function of the Fermi level at the chemical potential conditions corresponding to (a) the Al-rich limit and (b) the N-rich limit. The range of the Fermi level is given by the valence band maximum, which is set to be 0 eV and the conduction band minimum. The slope corresponds to charge state q. The formation energies of CeAl-ViAl (i = a, b, or c) and CeAl-VjN ( or ) defect complexes show similar behaviors against the Fermi level, respectively and therefore the most energetically preferable cases of i = c and j = c are plotted in these figures. The relaxed structure of the neutral CeAl defect viewed along the (c) and (d) [0001] directions, where the Ce atom is located at the Al site and the neighboring N atoms show atomic shifts away from the N sites (arrowheads).

It is noted that the Al-N bond length (1.89 Å) is ~29% shorter than that of Ce-N (2.44 Å) in the compound CeN. Hence, it may appear to be impossible to accommodate an isolated Ce atom into the w-AlN lattice. However, in the neutral CeAl defect structure, the tetrahedrally coordinated N atoms shift their positions and the resultant Ce-N bond length (2.22 Å) is ~17% longer than that of Al-N. We note here that the theoretical formation energy of a neutral CeAl defect is comparatively high (~3 eV) and so the doping of Ce atoms into the w-AlN lattice would seem to be very unlikely under ambient conditions. However, the present high pressure and high temperature extreme conditions are able to overcome the high formation energy of these CeAl luminescent centers.

In summary, we have synthesized Ce-doped w-AlN single crystals with high efficiency pink-colored luminescent centers through an optimized high pressure, high temperature flux method. The atomic structure of the luminescent center in w-AlN has been determined through direct observation of atomic-resolution ADF STEM imaging combined with systematic first-principles calculations. Our findings indicate that the present synthesis method enables the stable doping of isolated single Ce atoms and could be extended to control the emission color by choosing appropriate dopant elements. Moreover, it opens a new alternative strategy for doping size-mismatched functional atoms into wide band-gap materials, overcoming the high formation energies.


Single crystal synthesis

The single crystals of w-AlN:Ce were synthesized by the temperature gradient crystal growth method at 6.5 GPa and 1600°C for 20 hours, using a modified belt-type high-pressure apparatus with a bore diameter of 30 mm (Ref. 16, 17). The source of AlN powder (Toyo Soda, Type F) placed at upper part in the sample chamber was dissolved into the molten solvent and w-AlN crystals were precipitated at the cooler bottom in the chamber via a spontaneous nucleation process. We used a mixture of Li3AlN2 and Ba3Al2N4 as a solvent (typically, 1:1 molar ratio). In order to suppress carbon and oxygen contamination, we prepared solvents at 900°C under dry nitrogen atmosphere and then high-purity CeF3 powder (Rare metallic Co. Ltd, 4N grade, 0.5 wt%) as a dopant source was also mixed under the same atmosphere. The fundamental procedure for the crystal growth is similar to that of cubic boron nitride (c-BN) single crystals, where the solvents are Li3BN2 and Ba3B2N4 (Ref. 16, 17). The yields of w-AlN crystals are essentially 100% because the source of AlN powder was entirely dissolved into the solvents and precipitated as recrystallized w-AlN during growth duration, though the size of the crystals depend upon the position in the Mo sample chamber. Typical weights of the source and the solvent are 0.15 g and 0.18 g.

Electron microscope experiments and simulations

The surface of the obtained single crystals was cleaned by hot aqua regia. To directly image Ce dopants in the w-AlN lattice, an important prerequisite is to prepare clean and electron-transparent thin TEM specimens (<10 nm). To avoid any surface damage, we did not use conventional Ar-ion thinning but gently crushed the single crystals in ethanol and dispersed them onto an amorphous carbon grid. To remove hydrocarbon contamination, the grid was baked at 160°C for 8 hours in a clean vacuum and then transferred into a microscope. The atomic-resolution ADF STEM images were acquired with an aberration corrected Nion UltraSTEM 200, operated at 200 kV. To avoid beam damage, atomic-resolution observations were performed under a low beam current condition (~9 pA). Under this condition, no significant change can be recognized even scanning more than 30 times over the same region. The image simulation was carried out using the frozen phonon model with a probe-forming aperture half-angle of 30 mrad, the half-angle of the ADF detector spanning 63 to 409 mrad and incorporating the experimentally-measured, non-uniform detector efficiency25.

XANES experiments

The XANES measurements were carried out using the total electron yield method at BL11A in KEK-PF, Tsukuba, Japan.

First-principles calculations

The calculations were performed using the projector augmented-wave method26 and the HSE06 hybrid functional20,21 as implemented in the VASP code27,28. The defects were modeled using 192-atom supercells and spin polarization was allowed for all the defect species and charge states. A plane-wave cutoff energy of 400 eV and the Γ-point only k-point sampling were used in the calculations. To correct finite-size effects of the supercells with charged defects, approximate third-order image charge corrections reported in Ref. 29 were applied in conjunction with electrostatic potential alignment suggested in Ref. 30.