Growth of Zn–N Co-Doped Ga₂O₃ Films by a New Scheme with Enhanced Optical Properties

Abstract

Gallium oxide (Ga₂O₃), as a wide-bandgap semiconductor material, is highly expected to find extensive applications in optoelectronic devices, high-power electronics, gas sensors, etc. However, the photoelectric properties of Ga₂O₃ still need to be improved before its devices become commercially viable. As is well known, doping is an effective method to modulate the various properties of semiconductor materials. In this study, Zn–N co-doped Ga₂O₃ films with various doping concentrations were grown in situ on sapphire substrates by atomic layer deposition (ALD) at 250 °C, followed by post-annealing at 900 °C. The post-annealed undoped Ga₂O₃ film showed a highly preferential orientation, whereas with the increase in Zn doping concentration, the preferential orientation of Ga₂O₃ films was deteriorated, turning it into an amorphous state. The surface roughness of the Ga₂O₃ thin films is largely affected by doping. As a result of post-annealing, the bandgaps of the Ga₂O₃ films can be modulated from 4.69 eV to 5.41 eV by controlling the Zn–N co-doping concentrations. When deposited under optimum conditions, high-quality Zn–N co-doped Ga₂O₃ films showed higher transmittance, a larger bandgap, and fewer defects compared with undoped ones.

1. Introduction

Gallium oxide (Ga₂O₃) has been considered for a wide spectrum of applications in transparent conducting oxides, deep ultraviolet photodetectors, and high-voltage power electronic devices [1,2,3,4,5], thanks to its fascinating features, such as an ultra-wide bandgap of 4.9 eV [6,7], a high breakdown voltage (8 MV/cm) [8,9], and a high Baliga’s figure of merit (BFOM) of around 3444 [10,11]. However, due to the facile presence of oxygen vacancies, Ga₂O₃ generally displays n-type semiconducting properties [12], which limits its practical applications.

Doping with selective elements is an effective method to control the electronic and optical properties of semiconductors, including Ga₂O₃. In recent years, a series of Ga₂O₃-based materials with excellent properties have been successfully prepared based on this strategy [13,14]. The conductivity of Ga₂O₃ can be adjusted by doping with Sn, Ge, or Si elements [15,16,17,18], and the optical bandgap of Ga₂O₃ semiconductors can be tuned by doping with Zn and Al [19,20]. Great efforts have been made to improve the results of p-type doping of Ga₂O₃-based materials [19,21,22,23]. Given the similar ionic radii of N³⁻ and O²⁻, nitrogen is considered one of the most promising dopants for regulating the electronic and optical properties of Ga₂O₃-based materials. However, effective high-concentration nitrogen doping is usually difficult to achieve in most studies. Notably, high-quality nitrogen-doped p-type Ga₂O₃ materials were successfully prepared by the thermal oxidation of gallium nitride (GaN) on sapphire substrates at 1000–1100 °C [24,25]. First-principles calculations showed that tunable optical properties and even p-type conductivity can be achieved through co-doping with Zn and N. However, it is very difficult to achieve co-doping in the experiment. So far, the influence of Zn and N co-doping on the physical properties of Ga₂O₃ films is still unknown experimentally. Therefore, it is rewarding to take on this challenge in this work.

Several deposition techniques have been used to prepare Ga₂O₃ films, such as molecular beam epitaxy (MBE) [26,27], pulsed laser deposition (PLD) [28,29], magnetron sputtering, metal–organic chemical vapor deposition (MOCVD) [30], and atomic layer deposition (ALD) [31,32]. Compared with other thin film deposition techniques, ALD offers many advantages, such as precise thickness control, excellent 3D conformality, wafer-scale uniformity, compatibility with the mass production of semiconductor integrated circuits, etc. These advantages position ALD among the most promising ultra-thin and conformal thin film growth technologies, enabling precise dopant concentration control. However, the low deposition temperature characteristic of ALD has led to amorphous or polycrystalline Ga₂O₃ films. Furthermore, systematic investigations into controlled doping—particularly co-doping—of Ga₂O₃ via ALD remain scarce, hindering the mechanistic understanding of dopant effects in these thin films. In this work, the Zn–N co-doped β-Ga₂O₃ films were achieved by ALD and the effects of the doping concentrations of Zn and N elements on the crystallinity, optical bandgap, surface morphology, and photoluminescence properties were systematically investigated. GaN/ZnO nano-laminated films were grown by combining the plasma-enhanced ALD (PEALD) of GaN and the thermal ALD of ZnO at 250 °C, and then the Zn–N co-doped Ga₂O₃ films were prepared by the thermal oxidation of the deposited GaN/ZnO films at 900 °C. The Zn–N co-doping concentration was tuned by adjusting the ratio of the ALD growth cycles of the GaN and ZnO monolayers.

2. Experimental Procedures

2.1. Sample Preparation

The depositions of the GaN/ZnO nano-laminated films were conducted on c-plane (0001) sapphire (alpha-Al₂O₃) wafers at 250 °C in an ALD (R200, Picosun, Espoo, Finland) reactor. Prior to the reactions, the wafers were sequentially rinsed with acetone, ethyl alcohol, and deionized (DI) water in an ultrasonicator for 10 min, and then dried with a N₂ gun before being placed into the reaction chamber. The Zn–N co-doped Ga₂O₃ films were fabricated by the thermal oxidation of the ALD-deposited GaN/ZnO nano-laminated films. GaN was grown using triethylgallium (TEGa) and an Ar/N₂/H₂ (1:3:6) mixture gas plasma as the Ga and N precursors, respectively. ZnO was grown from the precursor of DI water and diethylzinc (DEZn) through thermal ALD. The precursor of TEGa was heated up to 60 °C, while the DI water and DEZn precursor was kept at room temperature (about 20 °C). N₂ (99.999%) was utilized as the carrier gas in the experiment. During the remote plasma process, the mixture plasma gas flow and RF power were set to 180 sccm and 2000 W, respectively. The PEALD growth sequence of GaN consisted of TEGa exposure (0.1 s), N₂ purging (5 s), Ar/N₂/H₂ plasma exposure (12 s), and N₂ purging (4 s). The thermal ALD growth sequence of ZnO consisted of DEZn exposure (0.1 s), N₂ purging (5 s), DI water exposure (0.1 s), and N₂ purging (5 s). After several cycles (n = 12, 6, or 3) of GaN deposition, 1 ZnO cycle was performed to complete a supercycle of Zn-doped GaN deposition. The Zn-doped GaN films were deposited after several supercycles. The details of the film composition are provided in

Table 1. Detailed processing parameters for deposition of Ga₂O₃-based thin film under different conditions.

Sample	Growth Cycles of GaN	Growth Cycles of ZnO	Number of Supercycle	Growth Cycles of Ga2O3
GaN	500	0	/	/
Ga:Zn = 12:1	468	39	39 (12GaN + 1ZnO)	/
Ga:Zn = 6:1	450	75	75 (6GaN + 1ZnO)	/
Ga:Zn = 3:1	375	125	125 (3GaN + 1ZnO)	/
Ga2O3	/	/	/	500

. In order to achieve Zn–N co-doped Ga₂O₃ and to activate the doping atoms, the as-grown Zn-doped GaN films were oxidized through post-annealing at 900 °C for 1 h, followed by naturally cooling down to room temperature. The resulting Zn–N co-doped Ga₂O₃ films are named as N-Ga(n)Zn(1), where n denotes the number of cycles of GaN in the supercycle of Zn-doped GaN deposition. For comparison, the GaN films without Zn doping were deposited after 500 cycles of GaN. The GaN films were transformed into N-doped Ga₂O₃ (N-Ga₂O₃) film by post-annealing at 900 °C in air. Furthermore, for comparison, a pure Ga₂O₃ film was deposited in situ using the following process parameters: TEG exposure (0.1 s), N₂ purging (6 s), O₂ plasma exposure (10 s), and N₂ purging (5 s).

2.2. Characterization

X-ray diffraction was employed to analyze the crystalline quality of the films (XRD, Bruker, Bremen, Germany) using Cu Kα radiation (45 kV, 40 mA, λ = 1.54056 Å) in the 2θ range of 20 to 120°. The morphology of the thin films was examined using atomic force microscopy (AFM, Bruker Dimension Icon, Bremen, Germany) in tapping mode. An ultraviolet–visible spectrophotometer was used to measure the absorption spectra of the thin films in the wavelength range of 190–800 nm. The Tauc plots based on the UV–vis data were further used to determine the optical bandgap. The microstructures of the thin films were investigated via transmission electron microscopy (TEM, JEOL JEM-F200, Tokyo, Japan). A secondary ion mass spectrometer (SIMS, Hesse, Germany) was employed to investigate the elemental distribution and diffusion in the films. The photoluminescence (PL) performance of the samples was evaluated by fluorescence spectrometry (FLS1000, Techcomp, Edinburgh, UK).

3. Results and Discussion

Figure 1 shows the XRD patterns of the Zn–N co-doped Ga₂O₃ films post-annealed at 900 °C with various doping concentrations. The undoped Ga₂O₃ film post-annealed at 900 °C shows four significant diffraction peaks, corresponding to the $\left(\overline{4}02\right)$, $\left(\overline{6}03\right)$, $\left(\overline{8}04\right)$, and $\left(\overline{10}05\right)$ peaks of β-Ga₂O₃. The N-doped Ga₂O₃ film (N-Ga₂O₃) exhibits a similar XRD diffraction pattern to the undoped Ga₂O₃ film, indicating that the pure GaN films are oxidized into N-doped Ga₂O₃ through post-annealing at 900 °C, and the nitrogen doping has little effect on the lattice structure and crystallinity of the Ga₂O₃ films. The above results reveal that both undoped and N-doped Ga₂O₃ films show a preferential orientation. However, with the increase in the doping concentration of the Zn element, the intensity of the X-ray diffraction peaks of the Ga₂O₃ films becomes weaker or even disappears, showing that Zn doping deteriorates the crystallinity of the Ga₂O₃ films. The degradation of the crystal structure at high Zn doping levels may be partly attributed to insufficient oxidation kinetics during the post-annealing of Zn-enriched GaN/ZnO nanolayers, resulting from inadequate oxygen partial pressure and annealing time at elevated temperatures.

The detailed microstructure of the post-annealed N-Ga(6)Zn(1) Ga₂O₃ film was analyzed using FE-TEM and the results are shown in Figure 2. As shown in Figure 2a–e, the films have an atomically sharp interface, and no buffer or nucleation layer is observed at the interface of Ga₂O₃/Al₂O₃. A well-aligned crystal lattice image can be seen in the Zn–N co-doped film of Ga₂O₃, without any metastable phases observed. The high-resolution TEM (HRTEM) images show a lattice spacing of 0.469 nm for Ga₂O₃, corresponding to the $\left(\overline{2}01\right)$ interplanar distance of the β-Ga₂O₃ structure. Combining the in-plane results of HRTEM and the out-of-plane results of XRD, the growth relationship can be summarized as β-Ga₂O₃ $\left(\overline{2}01\right)$//α-Al₂O₃ $\left(0001\right)$. Since the crystal planes of the β-Ga₂O₃ $\left(\overline{2}01\right)$ family are polar planes (i.e., consisting of only one type of atom, either Ga or O), it can be deduced that the self-limited alternating reactions process of ALD favors the formation of polar planes in the thin films [31]. The low magnification cross-sectional TEM imaging and the corresponding EDS mapping analyses were further carried out on the N-Ga(6)Zn(1) Ga₂O₃ film. As shown in Figure 2b–e, the EDS mapping of Ga reveals a strong and sharp contrast in the Ga₂O₃ film, while the contrast for the Zn element is obviously weak and uniformly distributed in the Ga₂O₃ film. Moreover, no evident diffusion of Ga into the sapphire substrate was observed, indicating that no solid reaction or interdiffusion occurred between Ga₂O₃ and Al₂O₃ during the high-temperature post-annealing process at 900 °C. The concentration of the oxygen element in the film edge is lower than in the substrate, which may be due to the presence of oxygen vacancies and the substitution of the N element. Because the doping concentration of nitrogen is very low and the nitrogen in the air is physically adsorbed onto the surface of the TEM sample, it is difficult to identify any N element present in the films by the EDS mapping results.

In order to confirm that nitrogen is indeed successfully incorporated into the Ga₂O₃ films, the vertical distribution of the N, Zn, and Ga elements in the films were further analyzed using SIMS. Figure 3 displays the SIMS results of the post-annealed N-Ga(12)Zn(1) film, where the presence of the Ga (blue), Zn (green), and N (red) elements are clearly observed by the SIMS elemental mapping, confirming the successful co-doping of Zn and N elements in the Ga₂O₃ film. It is worth noting that the N and Zn doping elements are uniformly distributed throughout the entire Ga₂O₃ film, indicating that the doping is very uniform.

Figure 4 shows the AFM images and the root mean square (RMS) average roughness of various samples. The surface morphology of the as-deposited pure GaN film exhibits a large number of dispersed particles, as seen in Figure 4(b1). Compared with the undoped GaN films, the surface morphology of the Zn-doped GaN films reveals much less or no gross grains (Figure 4(c1–e1)), with a much-reduced average roughness (Ra), indicating that Zn doping can improve the surface smoothness of the GaN films. As seen in Figure 4(a1), the in situ grown Ga₂O₃ film exhibits a comparable morphology quality with the Zn-doped GaN samples. After post-annealing at 900 °C, the pure GaN and Zn-doped GaN films transformed into the N-doped Ga₂O₃ thin film and Zn–N co-doped Ga₂O₃ thin films, respectively, and their AFM surface morphology images are shown in Figure 4(b2–e2). As shown in Figure 4f, the roughness of the Zn-doped GaN films is almost the same as the undoped Ga₂O₃ film. However, the average roughness of the undoped GaN film increases from 1.42 nm to 4.27 nm after post-annealing at 900 °C, with the particles merged together to form larger aggregates, as presented in Figure 4(b2). This could be attributed to the oxidation and the transformation from an amorphous to a crystalline state. The surface roughness of the Zn–N co-doped Ga₂O₃ films monotonically decreases with the increase in Zn doping concentration. More importantly, the surface morphology of the N-Ga(3)Zn(1) film is smoother than the undoped Ga₂O₃ film, which may be attributed to the amorphous microstructure as confirmed by the XRD results.

In order to characterize the optical properties and bandgap of the films, optical transmittance was measured using UV–vis spectroscopy in the wavelength region from 190 nm to 800 nm. As shown in Figure 5a, all of the as-deposited films exhibit outstanding optical transmittance in the visible wavelength range from 400 nm to 780 nm. The transmittance of all of the as-deposited films drops sharply in the UV spectral range of 300 to 190 nm. The transmittance of the as-grown GaN film is the lowest, indicating the strongest absorption of UV light. The as-grown N-Ga(12)Zn(1) film exhibits a similar transmittance to the as-grown GaN film, owing to the close chemical composition between them. The UV spectral transmittance increases noticeably when the Zn concentration increases, indicating that the optical bandgap increases with the increasing concentration of Zn. After post-annealing at 900 °C, the UV region transmittances of all of the films are improved and the absorption edges are obviously blue shifted. At the same time, the N-doped Ga₂O₃ film (transformed from the as-grown GaN film) has a similar transparency to the undoped Ga₂O₃ film, which further indicates that the GaN film has been successfully oxidized into N-doped Ga₂O₃ films. More importantly, the increase in the doping concentrations of Zn elements in the Zn–N co-doped Ga₂O₃ films causes the absorption band edge of the co-doped Ga₂O₃ to obviously blue shift, which is due to the rise in carrier concentration [33].

For a direct bandgap semiconductor material, we can obtain the optical bandgap (E_g) of the films by extrapolating the linear part of the Tauc plot to zero. The Tauc plot is the plot of ${\left(\alpha hv\right)}^2$ versus $h\upsilon$, and it can be described by the following equation:

$${\left(\alpha hv\right)}^2\propto B(hv-E_g),$$

(1)

where h is the Planck’s constant, $\alpha$ is the absorption coefficient, ν is the frequency of the incident light, B is a constant, and E_g is the bandgap energy.

As shown in Figure 5e, before post-annealing at 900 °C, the undoped Ga₂O₃ film has the largest bandgap of 4.69 eV, consistent with the previous reports [34]. It is worth noting that the bandgap of the deposited GaN film is 4.08 eV, higher than the reported 3.4 eV, which could be due to the presence of oxygen impurities [35]. The bandgap increases with the increase in ZnO/GaN deposition cycle ratio.

After post-annealing at 900 °C, the optical bandgaps of all of the films were increased, which is due to the fact that the GaN films were completely oxidized to N-doped Ga₂O₃ films and the defect density is reduced. The calculated bandgaps of the undoped Ga₂O₃, N-Ga₂O₃, N-Ga(12)Zn(1), N-Ga(6)Zn(1), and N-Ga(3)Zn(1) films after post-annealing are 5.03, 5.05 eV, 5.22 eV, 5.27 eV and 5.31 eV, respectively. In comparison to the undoped Ga₂O₃ film, the N- and Zn-doped samples have larger bandgaps. At the same time, the bandgap is enlarged with the increasing Zn concentration, which may be attributed to the Burstein–Moss effect [19,36].

The band structures and defect properties were investigated by photoluminescence spectroscopy, as demonstrated in Figure 6. Each of the samples shows a UV emission center around 327 nm, which can be attributed to the self-trapped excitons (STE). The blue emission peaks centered at 445 nm and 473 nm are considered to originate from the recombination of a trapped hole in an acceptor site with a trapped electron in a donor site, while the green emission peaks at 511 nm and 567 nm are caused by the neutral oxygen interstitial defects [37,38,39,40]. For the GaN, N-Ga(6)Zn(1), and Ga₂O₃ films, the intensity of the UV emission peak is strengthened after post-annealing, indicating that the number of non-radiative recombination centers are reduced [41]. As the numbers of donor defects and acceptor defects are suppressed after post-annealing, the blue emission of undoped Ga₂O₃ film is weakened [42]. However, the incorporation of N and Zn introduces an acceptor band into the N-Ga₂O₃ and N-Ga(12)Zn(1) films, and thereby the intensity of the blue emission peaks mounts up dramatically [23,43]. When the Zn doping concentration is raised further in the N-Ga(6)Zn(1) and N-Ga(3)Zn(1) films, the blue emission is suppressed due to the effect of concentration quenching [44]. After post-annealing, the intensity of green emission is reduced for all of the samples, suggesting that the quality of the films has been significantly enhanced.

4. Conclusions

In conclusion, Zn–N co-doped Ga₂O₃ films have been successfully grown on sapphire substrates using ALD followed with a post-annealing process. After being post-annealed at 900 °C, all of the films transform into a highly preferentially oriented β-Ga₂O₃, except for the sample of N-Ga(3)Zn(1), which remains amorphous. The HRTEM results reveal that an atomically sharp interface is formed and partial epitaxial growth takes place at the N-Ga(6)Zn(1) β-Ga₂O₃/α-Al₂O₃ interface. As the Zn doping concentration increases, the surface roughness of the films was reduced from 4.27 nm to 0.154 nm. In the visible wavelength region, all of the films exhibit an ultra-high optical transmittance, and the optical transmittance increases with the Zn doping concentration in the UV wavelength region. The bandgap of the Ga₂O₃ films can be adjusted from 5.07 eV to 5.41 eV by changing the Zn doping concentration. The PL spectra of the films exhibit several emission centers around 327 nm, 445 nm, 473 nm, 511 nm, and 567 nm under the excitation wavelength of 280 nm. A decline in the green luminescence was observed after post-annealing, demonstrating a decrease in oxygen interstitial defects. This work demonstrates that the structural and optical properties of β-Ga₂O₃ film can be well tuned by Zn–N co-doping and post-annealing. These findings facilitate the development of Ga₂O₃-based solar-blind ultraviolet photodetectors and other optoelectronic devices.