*3.3. Gallium Oxide Doping Issues and Recent Progress*

*β*-Ga2O<sup>3</sup> is very easily doped *n*-type to the degenerate state, *n*-type doped *β*-Ga2O<sup>3</sup> with carrier concentration from 10<sup>16</sup> to 10<sup>20</sup> cm−<sup>3</sup> [110,111] has been achieved by Sn and Ge doping by MBE, Si and Sn doping by MOVPE, and Sn doping by MOCVD [69]. A high mobility at room temperature of 145–184 cm2V −1 s −1 [100,101,112] has been reached by Si doping, and even till 10<sup>4</sup> cm2V −1 s <sup>−</sup><sup>1</sup> at 46 K [109]. Having a high critical field (5.2 MV.cm−<sup>1</sup> without intentional doping [113]), the *β*-Ga2O3devices demonstrate high performance. Nevertheless, all the Ga2O3devices demonstrated thus far have been unipolar in nature (i.e., only *n*-type). In order to realize the full potential for WBG opto-electronics *β*-Ga2O3and to sustain high breakdown voltage (>6.5 kV), we need vertical geometry bipolar-junctionbased devices. Therefore, the realization of *p*-type *β*-Ga2O<sup>3</sup> is a primary challenge today for the gallium oxide scientific community (Figure 4).

There is a tendency in oxide compounds to have *n*-type conductivity, caused by vacancies in the oxygen atoms. This, as well as the fact that it is a UWBG material, intrinsic conduction is rare and even causes *p*- and *n*-type doping tends not to be symmetrical. This asymmetry is seen in gallium oxide, the hole conductivity is poor and is likely the main limitation for development of gallium oxide technology. Fundamental restrictions such as this area recurring issue in oxides, such as: (i) acceptor point defects with high formation energy; (ii) native donor defects with low energy—resting holes; and (iii) *p*-type oxides suffer from a high effective mass of the holes (this results in a low mobility), due to the top of the VB predominantly from localized O 2-p derived orbits.

Native *p*-type conductivity: Using thermodynamical calculations for the point defects on gallium oxide it can be seen that gallium oxide is "lucky", as when *β*-Ga2O3is at <sup>500</sup> ◦C, *<sup>P</sup>hole* <sup>≈</sup> 1.33 <sup>×</sup> <sup>10</sup>−<sup>2</sup> atm with a hole concentration around *<sup>p</sup>* <sup>≈</sup> <sup>10</sup><sup>15</sup> cm−<sup>3</sup> [114]. Comparing this to calculations for ZnO gives *<sup>P</sup>hole* <sup>≈</sup>10<sup>3</sup> atm, for the same temperature. This divergence is believed to be from higher formation energy of the donor vacancies in *β*-Ga2O<sup>3</sup> (approximately 1 eV higher per vacancy), making compensation mechanism by point defects less favorable in gallium oxide than in ZnO. As a consequence, it can be expected that *p*-type samples of *β*-Ga2O<sup>3</sup> with higher carrier concentrations (then intrinsic) can be obtained when doping with shallow acceptor impurities.

The native hole concentration was investigated by Nanovation (SME, France) [114] where undoped *β*-Ga2O3thin film grown on c-sapphire substrates by pulsed laser deposition (PLD) showing resistivity of *<sup>ρ</sup>* = 1.8 <sup>×</sup> <sup>10</sup><sup>2</sup> <sup>Ω</sup>.cm, hole concentration of *<sup>p</sup>* = 2 <sup>×</sup>10<sup>13</sup> cm−<sup>3</sup> and a hole mobility of 4.2 cm2V −1 s −1 [114]. The determination of conductivity mechanism showed that Ga vacancies act as deep level acceptors with the activation energy of 0.56 eV in the low compensated sample, having *Ea* = 1.2 eV ionization energy. Later, the improvement was shown that native *p*-type conductivity by post-annealing in an oxygen atmosphere for *β*-Ga2O<sup>3</sup> thin film was grown on c-sapphire substrates by MOCVD [115]. After oxygen annealing, the hole concentration was increased from 5.6 <sup>×</sup> <sup>10</sup><sup>14</sup> cm−<sup>3</sup> to 5.6 <sup>×</sup> <sup>10</sup><sup>17</sup> cm−<sup>3</sup> at 850 K. The author claimed that the annealing effect is related to the formation of VGa —V<sup>O</sup> ++ complexes as a shallow acceptor center with *E<sup>a</sup>* = 0.17 eV activation energy.

Device applications require higher hole concentrations (at operating temperature), which could be achieved via external acceptor impurity incorporation.

There are already extensive theoretical studies (standard density functional theory (DFT and DFT with GGA+U) of acceptor impurity doping of *β*-Ga2O<sup>3</sup> in order to identify efficient *p*-type dopant. Kyrtsos et al. [116] demonstrated by DFT calculations that dopants, such as Zn, Li, and Mg, will introduce deep acceptor level with ionization energies of more than 1 eV, thus, they cannot contribute to the *p*-type conductivity. However, this result could be influenced by the underestimation of the bandgap due to the semi-local approach. Varley et al. [117] predicted that self-trapped holes are more favorable than delocalized holes due to their energies and by theoretical calculation (self-trapping energy is 0.53 eV and barrier to trapping is 0.10 eV). This indicates that free holes are unstable and will spontaneously localize towards small polarons.

Lyons [118] examined the elements of group 5 and group 12 (Be, Mg, Ca, Sr, Zn, Cd) as acceptor impurities in *β*-Ga2O<sup>3</sup> by hybrid DFT, all of them will exhibit the acceptor ionization levels of more than 1.3 eV. Mg was determined to be the most stable acceptor species, followed by Be. Sun et al. [119] used ab initio calculations to simulate the doping by Ge, Sn, Si, N, and Cl. Among them, N has been predicted to be a deep acceptor with an impurity level of 1.45 eV, as it has a similar atomic size as oxygen but has one less valence electron, and a higher *2p* orbital than oxygen. While all others act as donors, another ab initio calculation also demonstrated that nitrogen doping could introduce an acceptor level at 1.33 eV above the VBM.

Very recently, Goyal et al. [120] simulated a growth-annealing-quench sequence for hydrogen-assisted Mg doping in Ga2O<sup>3</sup> by using the first principles defect theory and defect equilibrium calculations. The H2O partial pressure and H exposure can strongly influence the Mg dopants concentration during the growth, by increasing the solubility limit of the acceptor, or by reducing the compensation. A conversion from *n*-type to *p*-type was achieved by annealing at O-rich/H-poor conditions. A Fermi level at +1.5 eV above the VB has been found after quenching.

Doping with two elements (co-doping) has been predicted by DFT which showed a promising method to obtain *p*-type *β*-Ga2O3, as it can break the solubility limit of monodoping and improves the photoelectric properties of semiconductor materials which results in increasing the conductivity.

The principle is to increase carrier concentration and decrease the compensating defect formation energy. This is inherently caused by the localized nature of the O2 *p*-derived VB that leads to difficulty in introducing shallow acceptors and large hole effective mass [121].

Co-doping has been successfully used for II-VI compounds, co-doping containing N (Zn-N, N-P, Al-N, and In-N) has been demonstrated to be an effective way to improve the *p*-type conductivity [122–124], in particular, Zhang et al. [124] predicted two shallow impurity levels above the VB of about 0.149 eV and 0.483 eV in N–Zn co-doped *β*-Ga2O3. Co-doping by N-P made an acceptor level decrease ~0.8 eV, and an impurity level appears at 0.55 eV above the VB of *β*-Ga2O3. A significant loss of holes' effective mass was also evidenced [124]. There are a few experimental works reported regarding *p*-type doping of gallium oxide. Mg-doped *β*-Ga2O3was studied by Qian et al. [125] for the photo-blind detector, and the *β*-Ga2O<sup>3</sup> containing 4.92 at% Mg has shown an acceptor level by XPS. A variation of bandgap has also been reported [83,126] however, the Hall effect measurement validity failed at room temperature due to the very high resistivity of the samples [127].

Suet al. [128] deposited Mg-Zn co-doped *β*-Ga2O<sup>3</sup> on sapphire (0001), however, antisites' impurity defects (i.e., ZnGa and GaZn) were determined as deep acceptors (0.79 eV for ZnGa and 1.00 eV for GaZn) by absorption spectra. Feng et al. [129] demonstrated Zn doping (1.3–3.6 at%) in *β*-Ga2O3nanowires can reduce the bandgap slightly, they also proved the *p*-type conductivity by making *p*-*n* junction. Chikoidze et al. [24] suggested that Zn in *β*-Ga2O<sup>3</sup> has an amphoteric nature: it can be an acceptor as ZnGa defect and at the same time, a donor being in Zn<sup>i</sup> interstitial sites. It was shown that in (0.5%) Zn:Ga2O<sup>3</sup> the auto-compensation of donor (Zn<sup>i</sup> ) -acceptor (ZnGa) defects takes place.

Islam et al. [130] reported that hydrogen annealing could vastly reduce the resistivity and reach a remarkable hole density of ~ 10<sup>15</sup> cm−<sup>3</sup> at room temperature. Besides, the ionization energy of acceptor is as low as 42 meV by incorporation of hydrogen in the lattice. This improvement is related to hydrogen decorated gallium vacancies VGa-H: during the diffusion of hydrogen into the Ga2O3crystal, H<sup>+</sup> absorbed at the surface will be attracted toward the VGa <sup>3</sup>−, it stabilizes the negative charge and thus lowers the acceptor level. This mechanism leads to H<sup>+</sup> decorated Ga-vacancy VGa-2H <sup>1</sup><sup>−</sup> and, therefore, the *p*-type conductivity.

Nitrogen-doped *p*-Ga2O<sup>3</sup> has been experimentally achieved by non-conventional growth technique. Wu et al. [131] demonstrated a multi-step structural phase transition growth from hexagonal P63mc GaN to rhombohedral R3C α-GaNxO3(1-x)/2 and realized the monolithic C2/m N-doped *β*-Ga2O<sup>3</sup> thin layer finally with an acceptor ionization

energy of 0.165 eV. The resistivity, hole concentration, and hole mobility are 17.0 Ω.cm, 1.56 <sup>×</sup> <sup>10</sup>16cm−<sup>3</sup> , and 23.6 cm2V −1 s −1 , respectively, by employing the Hall effect measurement. A performant field-effect transistor was also fabricated based on this *p*-type *β*-Ga2O3. Clearly, further experimental studies of optimal acceptor defects with room temperature activation are required.
