The hot-wall metal-organic chemical vapor deposition (MOCVD), previously shown to enable superior III-nitride material quality and high performance devices, has been explored for Mg doping of GaN. We have investigated the Mg incorporation in a wide doping range (2.45×1018 cm−3 up to 1.10×1020 cm−3) and demonstrate GaN:Mg with low background impurity concentrations under optimized growth conditions. Dopant and impurity levels are discussed in view of Ga supersaturation, which provides a unified concept to explain the complexity of growth conditions impact on Mg acceptor incorporation and compensation. The results are analyzed in relation to the extended defects, revealed by scanning transmission electron microscopy, x-ray diffraction, and surface morphology, and in correlation with the electrical properties obtained by Hall effect and capacitance–voltage (C–V) measurements. This allows to establish a comprehensive picture of GaN:Mg growth by hot-wall MOCVD providing guidance for growth parameters optimization depending on the targeted application. We show that substantially lower H concentration as compared to Mg acceptors can be achieved in GaN:Mg without any in situ or post-growth annealing resulting in p-type conductivity in as-grown material. State-of-the-art p-GaN layers with a low resistivity and a high free-hole density (0.77 Ω cm and 8.4×1017 cm3, respectively) are obtained after post-growth annealing demonstrating the viability of hot-wall MOCVD for growth of power electronic device structures.

Mg doping of GaN is essential for obtaining epitaxial layers with p-type conductivity for GaN-based devices such as light-emitting diodes (LEDs), normally-off high electron mobility transistors (HEMTs), and p-n diodes. Metal-organic chemical vapor deposition (MOCVD) is the technique of choice to grow large-scale GaN-based device structures and p-type doping in MOCVD GaN:Mg with free-hole concentrations in the range of 10161018 cm3 has successfully been demonstrated.1–3 

As-grown MOCVD Mg-doped GaN layers have limited or no p-type conductivity due to the formation of Mg–H complexes.4 Originating from the pyrolysis of hydrogen carrier gas and from the ammonia precursor cracking, atomic H is abundant in the MOCVD processes. The incorporated hydrogen must dissociate from the Mg–H complexes and out-diffuse to render GaN p-type, which requires post-growth treatment of the as-grown material.5,6 The activation is typically achieved via ex situ rapid or regular thermal annealing under N2 atmosphere. Due to the large Mg acceptor ionization energy, the free-hole concentration at room temperature is much lower than the Mg doping concentration. For this reason, high Mg concentration should be incorporated in GaN to reduce the ionization energy.1,7 However, above a certain level, typically in the low 1019 cm3 range, a further increase in Mg concentration leads to a reduction in free-hole concentrations.1,3,8 The latter is often associated with the formation of pyramidal inversion domains (PIDs).2,9 It has been shown that Mg atoms segregate at the PID (0001) boundaries rendering Mg atoms electrically inactive.3 Other works have suggested that the nitrogen vacancy, VN, and its complexes (e.g., Mg-VN) play a role in the compensation of Mg acceptors at high dopant concentrations.10,11 In addition, any donors, e.g., unintentional impurities such as Si and O or native defects such as Ga interstitials, will interfere negatively with the desired p-type conductivity. Notably, C is a major impurity in MOCVD due to the metal-organic precursor molecules cracking. The roles of C as a trap and compensation defect in GaN:Mg have been extensively discussed. Density functional theory calculations showed that C occupying Ga site is a shallow donor and it has the lowest formation energy for Fermi level positions close to the valence band maximum.12 Experimentally, it was demonstrated that the C incorporation leads to an increased donor concentration and a reduction in hole mobility, which was attributed to a donor-like +1/0 state of CN.3 Variation of growth conditions, such as V/III ratio, precursor, and gas flows, is expected to affect impurity incorporation and defect formation energies and, hence, their densities might influence the degree of passivation/compensation of Mg acceptors. Therefore, optimization and tuning of the growth process is explicitly necessary for acquiring low impurity and native defect levels in GaN:Mg. In addition, a characteristic delay in Mg incorporation relative to the onset of Cp2Mg precursor flow is common to MOCVD of Mg-doped GaN. Various measures including pre-flow of Cp2Mg to the reactor and purposeful design of the heat zone in the reactor13 have been adopted.

Continuous efforts in further improving p-type conductivity and free-hole properties, including both growth and post-growth acceptor activation, are undertaken to further the progress in GaN-based device performance.3,14 Despite the numerous investigations, no study exists on the capability of hot-wall MOCVD to deliver high-quality GaN:Mg layers with p-type conductivity. Because of its concept and the range of high deposition temperatures (up to 1400–1600 °C) achievable, the hot-wall MOCVD has successfully complied to conditions required for deposition of epitaxial layers of Al(Ga)N of semiconductor quality15,16 and reflected in demonstrating exciton luminescence in photoluminescence measurements17 and Mg-doped Al0.85Ga0.15N layers with low resistivity at room temperature.18 Recently, a room-temperature mobility above 2200 cm2/Vs of two-dimensional electron gas (2DEG) AlGaN/GaN HEMT heterostructures,19,20 a GaNSiC hybrid material for high-frequency and power electronics,21 and state-of-the-art N-polar AlN epitaxial layers22 using hot-wall MOCVD have been demonstrated. The hot-wall MOCVD concept enables reduced temperature gradients in both vertical and horizontal directions. It further allows for independent control of the gas phase chemistry over the substrate and the growth species surface diffusion, which may be exploited to control defect formation and impurity incorporation in a wider growth window.

In this work, we report a systematic study of Mg-doped GaN epitaxial layers grown by hot-wall MOCVD. A detailed investigation of the effects of growth and doping conditions on the incorporation of Mg, H, and C impurities is presented and discussed in terms of Ga supersaturation. The results are evaluated in relation to extended defects, revealed by transmission electron microscopy as well as x-ray diffraction and surface morphology, and correlated with the electrical and free-hole properties of the GaN:Mg layers obtained by Hall-effect and capacitance–voltage measurements. As a result, a comprehensive picture of hot-wall MOCVD of GaN:Mg is established, demonstrating the capabilities of this technique to deliver state-of-the-art p-type GaN.

A hot-wall MOCVD reactor in horizontal configuration was utilized for the epitaxial growth of all layers. Chemical–mechanical polished 4H-SiC substrates with on-axis (0001) orientation were used after wet chemical cleaning treatment. Before growth, the cleaned substrates were annealed and etched with hydrogen at 1340 °C in the MOCVD reactor. AlN nucleation layers (NLs) with a thickness of 50 nm were grown at 1250 °C with a V/III ratio of 1258, followed by growth of 500 nm-thick Mg-doped GaN layers. The growth process was performed under a constant ammonia (NH3) flow (2 l/min), a mixture of N2 and H2 as carrier gases, and a constant pressure of 100 mbar. NH3, trimethylaluminum (TMAl), trimethylgallium (TMGa), and bis(cyclopentadienyl)magnesium (Cp2Mg) were used as the N, Al, Ga, and Mg precursors, respectively. The growth conditions of the GaN:Mg layers are summarized in Table I. The unintentionally doped (UID) GaN reference sample and the Mg-doped GaN samples with Mg doping concentration ranging from 2.45×1018 cm−3 up to 1.10×1020 cm−3 are denoted as M0 and M1–M9, respectively. Several sets of Mg-doped GaN layers were grown by systematically varying one of the following growth parameters: (i) Cp2Mg/TMGa ratio—samples M1–M5 grown under optimized conditions with low carrier gas flows, where the Cp2Mg/TMGa is varied between 0.033% and 0.5% and samples M7 and M8 grown with high carrier gas flows with Cp2Mg/TMG of 0.335% and 0.5%, respectively; (ii) V/III ratio—samples M4, M6, and M7 grown with Cp2Mg/TMGa of 0.335% and V/III ratios of 906, 1811, and 453, respectively; and (iii) growth temperature (Tg)—samples M8 and M9 with Tg=1120°C and Tg=1040°C, respectively. In addition, a sample, Mopt was grown on 1 μm-thick undoped GaN using the conditions of M3. The Mg acceptors were activated using optimized ex situ annealing process at 900 °C in N2 atmosphere in a rapid thermal processing equipment.

TABLE I.

A summary of the growth conditions for the GaN:Mg layers studied in this work: growth temperature (Tg), growth pressure, V/III ratio, growth rate, Cp2Mg/TMGa ratio, and H2 and N2 carrier gas flows.

SampleTg (°C)Pressure (mbar)V/III ratioGrowth rate (μm/h)Cp2Mg/TMGa (%)H2 (l/min)N2 (l/min)
M0 1120 100 906 0.70 19 
M1 1120 100 906 0.70 0.033 19 
M2 1120 100 906 0.60 0.082 19 
M3 1120 100 906 0.70 0.167 19 
M4 1120 100 906 0.70 0.335 19 
M5 1120 100 906 0.60 0.500 19 
M6 1120 100 1811 0.14 0.335 19 
M7 1120 100 453 1.30 0.335 25 12 
M8 1120 100 453 1.06 0.500 25 12 
M9 1040 100 453 1.20 0.500 25 12 
Mopt 1120 100 906 0.60 0.167 19 
SampleTg (°C)Pressure (mbar)V/III ratioGrowth rate (μm/h)Cp2Mg/TMGa (%)H2 (l/min)N2 (l/min)
M0 1120 100 906 0.70 19 
M1 1120 100 906 0.70 0.033 19 
M2 1120 100 906 0.60 0.082 19 
M3 1120 100 906 0.70 0.167 19 
M4 1120 100 906 0.70 0.335 19 
M5 1120 100 906 0.60 0.500 19 
M6 1120 100 1811 0.14 0.335 19 
M7 1120 100 453 1.30 0.335 25 12 
M8 1120 100 453 1.06 0.500 25 12 
M9 1040 100 453 1.20 0.500 25 12 
Mopt 1120 100 906 0.60 0.167 19 

The Mg and the background impurity (H, C, Si, O) depth profiles in as-grown and annealed samples were measured by secondary ion mass spectroscopy (SIMS). The detection limit was 1×1016 cm−3 for Mg, 1×1017 cm−3 for H, (3–5)×1015 cm3 for Si and O, and 1×1015 cm−3 for C. The average concentrations for all the elements were estimated using the SIMSview program from EAG Labs.23 The surface morphology of the layers was studied by atomic force microscopy (AFM) using a Veeco Dimension 3100 scanning probe microscope in tapping mode. Both 5×5μm2 and 80×80μm2 images were acquired. Spectroscopic ellipsometry measurements were performed on a J. A. Woollam RC2-XI ellipsometer in the ultraviolet-visible spectral range (0.7–5.9 eV) for the determination of the layer thicknesses. The crystalline quality of the layers was evaluated by high-resolution x-ray diffraction (HRXRD) using a PANalytical Empyrean diffractometer. A hybrid monochromator consisting of a parabolic x-ray mirror and a two-bounce Ge(220) crystal resulting in CuKα1 radiation with wavelength λ = 1.540 597 4 Å at the incident x-ray beam and a symmetric three-bounce Ge(220) analyzer on the detector side was used. HRXRD 2θω scans of the symmetric 0006 and the asymmetric 101¯5 Bragg peaks were used for the determination of the c and a lattice parameters, respectively. The screw- and edge-type dislocation densities, NS and NE, were estimated using the tilt and twist angles, αS and αE, respectively,24,25 and the lattice parameters of a bulk GaN,26 

NS=αS24.35bc2,NE=αE24.35ba2.
(1)

The tilt angle was determined by the Williamson–Hall plots of the symmetric 0002, 0004, and 0006 diffraction peaks of GaN,25 and the magnitude of the Burger vector along the c-axis (bc=0.5185 nm) was used for an estimation of the screw-type dislocation density. For the edge dislocation density estimation, a method is proposed by Srikant et al.,24 where the full-width at half-maximum (FWHM) of the rocking curves from the asymmetric 101¯1, 101¯2, 101¯3, 101¯4101¯5, and 303¯2 diffraction peaks and the Burger vector along the a-axis (ba=0.3189 nm) are used.

The structural quality of selected samples was further assessed by transmission electron microscopy (TEM) using the double corrected Linköping FEI Titan3 60-300 microscope, operated in scanning TEM (STEM) mode and 300 kV. Images were acquired using annular dark-field (ADF) and annular bright-field (ABF) detectors (collection angles 21–200 and 4–43 mrad, respectively), revealing both atomic number and strain contrast. Electron energy loss spectroscopy (EELS) thickness measurements were performed by acquisition using a Gatan GIF Quantum detector with 10 ms exposure time, 0.25 eV/channel dispersion (50.0 to 462.0 eV), and convergence and collection angles of 22.0 and 15.0 mrad, respectively. Calculations were done in the embedded function in Gatan Digital Micrograph software. Samples were mechanically cut and mounted in Ti grids, followed by mechanical thinning and Ar-ion milling (Gatan PIPS Model 691). This produced two projection directions: [11¯00] and [112¯0].

The free-hole concentration and mobility of the as-grown and annealed samples were determined by Hall-effect measurements performed at room temperature in Van der Pauw configuration using a Linseis HCS1 instrument. For this purpose, Ni/Au (5 nm/250 nm) ohmic contacts were deposited by thermal evaporation and annealed at 450 °C in air. Note that the Hall-effect measurements of as-grown and annealed samples were performed on the same piece for a given growth condition. The effective dopant concentration NAND in samples M1–M8 was extracted by capacitance–voltage (C–V) measurements, using a Hg-probe setup with a 4284A LCR meter from Agilent. The C–V data were acquired in series measurement mode in the 1–10 kHz frequency range.

The concentrations of Mg dopants, H and C impurity, the root mean square surface roughness, the density of screw and edge type dislocations, and the free-hole concentration in as-grown GaN and the free-hole concentration, mobility, and resistivity of the respective thermally activated GaN samples are summarized in Table II.

TABLE II.

Concentrations of Mg dopants ([Mg]), H ([H]) and C ([C]) unintentional impurities, root mean square (RMS) surface roughness, screw (NS) and edge (NE) dislocation densities, hole concentration (p0) in the as-grown GaN:Mg layers, and hole concentration (p), mobility (μ), and resistivity (ρ) in the respective layers after annealing.

As-grownAnnealed
Sample[Mg] (cm−3)[H] (cm−3)[C] (cm−3)RMS (nm)NS (cm−2)NE (cm−2)p0 (1016cm−3)p (1017cm−3)μ (cm2/V.s)ρ (Ω.cm)
M0 1.5 × 1017 9.9 × 1015 0.15 1.7 × 107 5.0 × 108 … … … … 
M1 2.5 × 1018 2.4 × 1018 9.8 × 1015 0.22 3.8 × 107 5.9 × 108 N/A N/A … … 
M2 6.0 × 1018 4.8 × 1018 6.0 × 1015 0.22 4.0 × 107 6.2 × 108 8.4 5.8 10 1.08 
M3 1.6 × 1019 1.5 × 1019 1.5 × 1016 0.19 2.1 × 107 5.4 × 108 0.76 6.5 1.08 
M4 2.4 × 1019 2.2 × 1019 1.5 × 1016 0.19 3.7 × 107 4.8 × 108 N/A 5.3 1.30 
M5 6.1 × 1019 9.6 × 1018 8 × 1015 0.26 2.9 × 107 8.1 × 108 9.6 2.3 2.90 
M6 1.6 × 1019 1.3 × 1019 5.3 × 1015 0.17 3.0 × 107 7.4 × 108 3.1 8.9 0.88 
M7 7.0 × 1019 1.4 × 1019 2.1 × 1016 0.20 2.4 × 107 9.0 × 108 0.61 0.9 7.78 
M8 6.7 × 1019 1.4 × 1019 2.3 × 1016 0.45 5.4 × 107 7.0 × 108 0.02 1.1 8.73 
M9 1.1 × 1020 2.9 × 1019 3.1 × 1017 1.00 … … N/A N/A … … 
Mopt 1.6 × 1019 1.5 × 1019 1.5 × 1016 0.18 1.6 × 107 4.3 × 108 1.67 8.4 10 0.77 
As-grownAnnealed
Sample[Mg] (cm−3)[H] (cm−3)[C] (cm−3)RMS (nm)NS (cm−2)NE (cm−2)p0 (1016cm−3)p (1017cm−3)μ (cm2/V.s)ρ (Ω.cm)
M0 1.5 × 1017 9.9 × 1015 0.15 1.7 × 107 5.0 × 108 … … … … 
M1 2.5 × 1018 2.4 × 1018 9.8 × 1015 0.22 3.8 × 107 5.9 × 108 N/A N/A … … 
M2 6.0 × 1018 4.8 × 1018 6.0 × 1015 0.22 4.0 × 107 6.2 × 108 8.4 5.8 10 1.08 
M3 1.6 × 1019 1.5 × 1019 1.5 × 1016 0.19 2.1 × 107 5.4 × 108 0.76 6.5 1.08 
M4 2.4 × 1019 2.2 × 1019 1.5 × 1016 0.19 3.7 × 107 4.8 × 108 N/A 5.3 1.30 
M5 6.1 × 1019 9.6 × 1018 8 × 1015 0.26 2.9 × 107 8.1 × 108 9.6 2.3 2.90 
M6 1.6 × 1019 1.3 × 1019 5.3 × 1015 0.17 3.0 × 107 7.4 × 108 3.1 8.9 0.88 
M7 7.0 × 1019 1.4 × 1019 2.1 × 1016 0.20 2.4 × 107 9.0 × 108 0.61 0.9 7.78 
M8 6.7 × 1019 1.4 × 1019 2.3 × 1016 0.45 5.4 × 107 7.0 × 108 0.02 1.1 8.73 
M9 1.1 × 1020 2.9 × 1019 3.1 × 1017 1.00 … … N/A N/A … … 
Mopt 1.6 × 1019 1.5 × 1019 1.5 × 1016 0.18 1.6 × 107 4.3 × 108 1.67 8.4 10 0.77 

Figure 1 shows the Mg and H concentrations in the GaN:Mg layers as a function of the ratio between the dopant precursor and the TMGa in the gas phase, Cp2Mg/TMGa. As expected, Cp2Mg/TMGa has a significant effect on the Mg incorporation. For samples M1–M5 grown under optimized growth conditions of V/III ratio of 906 and growth rate of 0.60–0.70 μm/h, [Mg] increases linearly with Cp2Mg/TMGa. A maximum Mg concentration of 6.0×1019 cm−3 at Cp2Mg/TMGa = 0.5% is reached. We deduce a dopant incorporation efficiency of 0.24 from the dependence of the atomic fraction of magnesium [Mg]/[Ga] in the lattice as a function of the Cp2/TMGa flux ratio in the vapor phase (see Fig. S1 in the supplementary material). The observed trends (Fig. 1 and Fig. S1 in the supplementary material) are consistent with the previously reported results for MOCVD GaN:Mg layers.2,8,27,28 We note that the characteristic delay in Mg incorporation, typical for MOCVD processes, is also observed in our experiments. The abruptness of the Mg doping profile can be influenced by increasing the hydrogen carrier flow as reported elsewhere.29 

FIG. 1.

Mg and H concentrations ([Mg] and [H]) determined by SIMS as a function of Cp2Mg/TMGa ratio in the as-grown GaN:Mg layers. Squares and triangles represent [Mg] and [H], respectively. The sample labels according to Table I are indicated in the figure. The blue line represents a linear fit to the [Mg] data points for samples M1–M5 grown under the same growth conditions.

FIG. 1.

Mg and H concentrations ([Mg] and [H]) determined by SIMS as a function of Cp2Mg/TMGa ratio in the as-grown GaN:Mg layers. Squares and triangles represent [Mg] and [H], respectively. The sample labels according to Table I are indicated in the figure. The blue line represents a linear fit to the [Mg] data points for samples M1–M5 grown under the same growth conditions.

Close modal

In order to explore the limitations in Mg incorporation levels, additional GaN:Mg layers, M6 - M9, with the two highest Cp2Mg/TMGa ratios of 0.335% and 0.5% but with different V/III ratios, temperatures, and growth rates were prepared. Samples M6 and M7 were grown with the same Cp2Mg/TMGa ratio of 0.335% as for sample M4 but with twice as high V/III ratio of 1811 and reduced by half V/III ratio of 453, respectively. It can be noted that for a constant CpMg/TMGa ratio, the increase in the V/III ratio leads to 1.5 times lower Mg concentration, bringing it below the threshold for PID formation (for details on PID, see Sec. III B). When the V/III is decreased to 453, a significant increase in [Mg] to 7.0×1019 cm3 is observed (M7). This is almost three times higher compared to the respective [Mg] in M4 with a V/III ratio of 906. Note that higher gas carrier flows of H2=25 and N2=12 l/min (Table I) were used for M7 but their ratio was kept the same as for the case of samples M0–M6. A detailed study on the effects of carrier gas flows and composition is reported elsewhere.29 In the next step, a GaN:Mg with the same growth parameters as for M7 but with a Cp2Mg/TMGa of 0.5% was grown (sample M8) to investigate whether Mg incorporation can be further boosted by increasing Cp2Mg/TMGa at these specific growth conditions. Interestingly, [Mg] in M8 is found to be 6.7×1019 cm−3, which is very similar to [Mg] in M7. This indicates that at low V/III ratios a saturation of Mg incorporation is already reached at a Cp2Mg/TMGa of 0.335%. The lower V/III ratio is associated with a higher growth rate, which can shift the balance from Mg atom substitutionally replacing a Ga atom toward remaining as an adatom providing potential explanation for the observed Mg saturation. Finally, a GaN:Mg with a V/III ratio of 453, Cp2Mg/TMGa ratio of 0.5% but at a lower temperature of 1040 °C—M9 was grown (the rest of the samples were all grown at 1120 °C). This also led to a slightly different growth rate of 1.20 μm/h compared to M8. The lower growth temperature resulted in high Mg incorporation level in the range of 1×1020 cm−3.

Within the medium-to-high Mg concentration range up to 2.4×1019 cm3, [H] follows closely [Mg] being slightly lower than the expected 1:1 ratio (see Fig. 1, Table II, and Fig. S2 in the supplementary material). Consequently, as-grown GaN:Mg exhibits p-type conductivity even before the annealing, as will be discussed later (see Sec. III C). The H levels in all GaN:Mg layers with [Mg] above 2.4×1019 cm−3 are below 3×1019 cm−3, being significantly lower than the respective [Mg] in agreement with earlier observations1,2 (see Fig. 1 and Fig. S2 in the supplementary material).

Our results indicate that the Mg incorporation, as expected, is directly related to the Cp2Mg/TMGa ratio, more specifically to the Mg precursor flow in the reactor. In addition, V/III ratio, growth rate, growth temperature, and pressure are impacting the growth process and, hence, incorporation of impurities and formation of native defects. It has been demonstrated that for MOCVD growth the concept of supersaturation, which measures the deviation from the thermodynamic equilibrium, is suitable to evaluate the impact of growth conditions in their complexity.30 Therefore, we present in Fig. 2 [Mg] as a function of the Ga supersaturation. The latter was calculated following Ref. 30 and the respective growth parameters in Table I. Notably, the supersaturation has a strong effect on the incorporated [Mg] for samples with a constant Cp2Mg/TMGa ratio. For instance, at Cp2Mg/TMGa of 0.335%, the decrease in supersaturation from 160 (M4) to 79 (M6) results in 65% decrease in [Mg] while increasing it to 184 (M7) leads to a steep rise of [Mg] by 165%. When the Ga supersaturation for constant Cp2Mg/TMGa of 0.5% increased from 184 (M8) to 1000 (M9) [Mg] is increased by 65%. For the first case (M4, M6, and M7), the increase in Ga supersaturation was achieved by varying the TMGa flow and by decreasing the growth temperature for the latter case (M8–M9). The ratio F (the ratio of the input mole fraction of H2 to the total amount of the two carrier gases, H2 and N2), which also strongly affects the supersaturation, is the same for all samples. These observations show that the Mg incorporation in GaN is a multi-dimensional process and the trends related to different growth parameters, e.g., V/III ratio, growth temperature, and carrier gas composition, can be assessed by taking into account only the Cp2Mg precursor flow and the Ga supersaturation. Clearly, Mg incorporation follows the same trend as Ga supersaturation, i.e., the available Ga species in the reaction. Since Mg atoms compete with Ga atoms for incorporation at the same lattice site, at a constant Ga supersaturation the balance of Mg/Ga atoms incorporation is shifted more toward the Mg side when the Cp2Mg precursor flow, i.e., the available Mg atoms in the reaction are increased.

FIG. 2.

Mg concentration ([Mg]) from Fig. 1 as a function of Ga supersaturation in the as-grown GaN:Mg layers. The sample labels according to Table I are indicated in the figure. The same as in Fig. 1, blue symbols indicate samples grown with 19 l of H2 and 9 l of N2 while green symbols indicate samples grown with 25 l of H2 and 12 l N2. The blue dashed arrow indicates the increase of Cp2Mg/TMGa at constant supersaturation, the red dashed arrows indicate variation in supersaturation due to change of TMGa flow at a constant Cp2Mg/TMGa of 0.335% and the green arrow indicates variation in supersaturation due to growth temperature at a constant Cp2Mg/TMGa of 0.5% for specific sets of growth conditions.

FIG. 2.

Mg concentration ([Mg]) from Fig. 1 as a function of Ga supersaturation in the as-grown GaN:Mg layers. The sample labels according to Table I are indicated in the figure. The same as in Fig. 1, blue symbols indicate samples grown with 19 l of H2 and 9 l of N2 while green symbols indicate samples grown with 25 l of H2 and 12 l N2. The blue dashed arrow indicates the increase of Cp2Mg/TMGa at constant supersaturation, the red dashed arrows indicate variation in supersaturation due to change of TMGa flow at a constant Cp2Mg/TMGa of 0.335% and the green arrow indicates variation in supersaturation due to growth temperature at a constant Cp2Mg/TMGa of 0.5% for specific sets of growth conditions.

Close modal

O and Si impurity levels in our GaN:Mg layers were found to be at the SIMS detection limit of (3–5)×1015 cm3 in all samples. Hence, only C from the unintentional impurities is further discussed in its potential role as a compensating donor and a trap affecting the free-hole properties. The respective C concentrations in the GaN:Mg films are presented in Fig. 3 as a function of Ga supersaturation. An overall linear increase in [C] with increasing Ga supersaturation is evident independently of whether the change in supersaturation is achieved via varying V/III ratio or growth temperature. The increase in [C] at constant supersaturation could be associated with an increase in the Cp2Mg/TMGa ratio.31 

FIG. 3.

Carbon concentration ([C]) in the as-grown GaN:Mg layers vs Ga supersaturation. The sample labels according to Table I are indicated in the figure. The same as in Fig. 1, blue symbol indicates samples grown with 19 l of H2 and 9 l of N2 while green symbols indicate samples grown with 25 l of H2 and 12 l of N2.

FIG. 3.

Carbon concentration ([C]) in the as-grown GaN:Mg layers vs Ga supersaturation. The sample labels according to Table I are indicated in the figure. The same as in Fig. 1, blue symbol indicates samples grown with 19 l of H2 and 9 l of N2 while green symbols indicate samples grown with 25 l of H2 and 12 l of N2.

Close modal

As can be seen from Table II, a small variation of the screw-type dislocations density upon increasing the Mg concentration is observed in the as-grown GaN:Mg layers. The NS remains very similar to the value found in the undoped GaN layer. On the other hand, NE increases from (5–6)×108 cm2 to (7–9)×108 cm2 for [Mg] 2.4×1019 cm3. For comparison, in Mg doped InN layers, the edge dislocation density remains similar with increasing Mg, while screw dislocation density increases accompanied by a polarity inversion.32 

STEM analysis was performed to provide further insight into the microstructure of the GaN:Mg layers. The results shown in Fig. 4 verify that extended defects are not generated in layers doped with Mg up to 1.6×1019 cm−3 [Figs. 4(a) and 4(b)]. For [Mg] = 2.4×1019 cm−3 (sample M4), sporadic PID defects are present [Figs. 4(c) and 4(e)]. Incorporating more Mg (sample M8 with [Mg]=6.7×1019) leads to a notable increase in PID density [Fig. 4(d)]. The average width of the PIDs in the latter case is estimated to be 5.4±0.9 nm [Fig. 4(d)], in range with the previously reported values.2,3 Polarity inversion inside the pyramids is confirmed by annular bright-field (ABF) STEM (see Fig. S3 in the supplementary material). STEM analysis further revealed segregation of Mg at the PID (0001) facet (Fig. S3 in the supplementary material) in concordance with earlier observations.2,3 Additionally, our comprehensive EELS analysis reveals that Mg atoms are segregated not only at the (0001) planes of the PIDs, but they are also found at their inclined facets (further details on the EELS analysis will be reported elsewhere).

FIG. 4.

STEM images of as-grown GaN:Mg layers with different [Mg]: (a) M0—UID, (b) M31.6×1019 cm3, (c) M42.4×1019 cm3, and (d) M86.7×1019 cm3. Arrows in (c) and (d) highlight some of the polarity inversion domains (PIDs). Higher magnification images of PIDs for M4 and M8 are shown in (e) and (f), respectively (both in the [11¯00] projection).

FIG. 4.

STEM images of as-grown GaN:Mg layers with different [Mg]: (a) M0—UID, (b) M31.6×1019 cm3, (c) M42.4×1019 cm3, and (d) M86.7×1019 cm3. Arrows in (c) and (d) highlight some of the polarity inversion domains (PIDs). Higher magnification images of PIDs for M4 and M8 are shown in (e) and (f), respectively (both in the [11¯00] projection).

Close modal

A crude estimate of the PID density results in a range between 1.9 and 8.9×1016 cm3 by sampling different regions of M8 [Fig. 4(d)]. The extended range originates from a thickness measurement by EELS acquired in an adjacent region. An estimate of the number of Mg atoms associated with the PIDs can be obtained by following Narita et al.,2,3 for calculating the Mg atoms bound to the top facet and adding the herein observed segregation of Mg to the side facets. Accordingly, [Mg] bound to the PIDs in M8 is estimated to be from 1.6×1019 to 7.7×1019 cm3. We expect the actual value to be on the higher end of the range since the PID density measurement was performed close to an edge where the measured thickness is on the lower end of our estimate. In comparison, the difference between [Mg] and [H] in this sample as measured by SIMS amounts to 5.3×1019 cm3 (Table II and Fig. S2 in the supplementary material). It is, therefore, inferred that the difference between [Mg] and [H] at doping levels above 2.4×1019 cm3 can be accounted for by Mg segregated at all the interfaces of the PIDs.

The presence of PIDs is also expected to affect the surface morphology, in particular, for the high-defect-density layers with [Mg] 2.4×1019 cm3. The UID GaN and GaN:Mg layers M0–M5 grown with the optimal growth conditions exhibit step-flow growth mode and the terrace size is similar for the undoped and doped layers up to [Mg] of 6.1×1019 cm3 as can be seen from the 5×5μm2 micrographs in Fig. 5. The root-mean-square (RMS) roughness for M0–M5 is in the range of 0.15–0.26 nm. However, on a larger scale, changes in morphology could be observed already at moderate Mg concentrations. Figure 6 shows representative 80×80μm2 micrographs of selected samples. It is seen that already for Mg doping of 1.6×1019 cm3 (M3) there is a slight increase in RMS as compared to M1. The RMS of sample M4, for which sporadic PIDs start to appear, is further increased by a factor of two (Fig. 6). With a further increase in Mg concentration, one can observe well defined hexagonal hillocks [Fig. 4(d)]. Their density and height is increasing with increasing [Mg], which is manifested in higher RMS surface roughness. This surface deterioration can be correlated with the enhanced PID density that takes place with increasing [Mg]. Previously, it was shown that for [Mg]  1.0×1020 cm−3 the formation of hexagonal hillocks and platelets can be associated with a polarity inversion to N-polar GaN.33–35 Our results indicate that this is a gradual process that starts already at [Mg] = 2.4×1019 cm−3 and is associated with the PID formation in the overall Ga-polar GaN. Note that this is not accompanied by any noticeable increase in screw dislocation densities as earlier suggested for the case of full polarity inversion in GaN33 and experimentally observed for InN.32 For example, samples M3 and M8 have very similar NS (Table II) while their PID density differs by orders of magnitude.

FIG. 5.

5×5μm2 AFM images of as-grown GaN:Mg layers with different [Mg]: M0—UID, M12.5×1018 cm−3, M26.0×1018 cm−3, M31.6×1019 cm−3, M42.4×1019 cm−3, and M56.1×1019 cm−3. For further details on the samples, see Tables I and II.

FIG. 5.

5×5μm2 AFM images of as-grown GaN:Mg layers with different [Mg]: M0—UID, M12.5×1018 cm−3, M26.0×1018 cm−3, M31.6×1019 cm−3, M42.4×1019 cm−3, and M56.1×1019 cm−3. For further details on the samples, see Tables I and II.

Close modal
FIG. 6.

Large-area 80×80μm2 AFM images of GaN layers doped with different Mg levels: M12.5×1018 cm−3, M31.6×1019 cm−3, M42.4×1019 cm−3, M56.10×1019 cm−3, M86.7×1019 cm−3, and M91.1×1020 cm−3. Note the different scale bar for images on the top and bottom rows. For further details on the samples, see Tables I and II.

FIG. 6.

Large-area 80×80μm2 AFM images of GaN layers doped with different Mg levels: M12.5×1018 cm−3, M31.6×1019 cm−3, M42.4×1019 cm−3, M56.10×1019 cm−3, M86.7×1019 cm−3, and M91.1×1020 cm−3. Note the different scale bar for images on the top and bottom rows. For further details on the samples, see Tables I and II.

Close modal

Figure 7 shows the net acceptor concentration NAND and the free-hole concentrations in the annealed samples. The relative difference between Mg and H concentrations with respect to the total Mg concentration, ([Mg]–[H])/[Mg] in the as-grown GaN:Mg is plotted as a function of [Mg] in Fig. 8. Assuming that all H binds to the available MgGa,36 this fraction represents the total amount of Mg atoms not bound to H and it may include: (i) MgGa acceptors not passivated by H and/or (ii) Mg incorporated at a different site in the GaN crystal lattice, e.g., Mg interstitial and Mg segregated on the PID facets. Two distinct behaviors of (NAND)/p and ([Mg]–[H])/[Mg] can be discerned below and above [Mg] = 2.4×1019 cm3, respectively.

FIG. 7.

Net acceptor concentration NAND (from C–V measurements) and free-hole concentrations (from Hall) in the annealed p-GaN as a function of [Mg]. The red and orange curves are guide to the eye. The blue solid line corresponds to NAND = [Mg].

FIG. 7.

Net acceptor concentration NAND (from C–V measurements) and free-hole concentrations (from Hall) in the annealed p-GaN as a function of [Mg]. The red and orange curves are guide to the eye. The blue solid line corresponds to NAND = [Mg].

Close modal
FIG. 8.

Fraction of non-passivated Mg ([Mg]–[H])/[Mg] in the as-grown GaN:Mg layers as a function of Mg doping.

FIG. 8.

Fraction of non-passivated Mg ([Mg]–[H])/[Mg] in the as-grown GaN:Mg layers as a function of Mg doping.

Close modal

Up to 2.4×1019 cm3, the NAND follows closely [Mg], indicating that Mg is incorporated as an acceptor within this doping range. A certain fraction between 4% and 21% of [Mg] remains non-passivated by H (Fig. 8 and Table II). Since all Mg atoms are incorporated as acceptors (Fig. 7), this result implies that [Mg]–[H])/[Mg] represents the fraction of Mg atoms of type (i), i.e., the isolated MgGa acceptor not passivated by H.37 As a result, as-grown GaN:Mg manifests a net p-type conductivity with free hole concentration in the low to high 1016 cm3 range (samples M2, M3, and M6, see Table II). These findings are very interesting since they are contrasted with previous reports, where [H] is equal to or larger than2,8 Mg and indicates the expanded opportunities offered by the hot-wall MOCVD concept. The higher the non-passivated Mg fraction, the higher the free-hole concentration in the as-grown GaN:Mg for this doping range. It has been shown theoretically that higher amounts of H during growth can be used to increase acceptor incorporation and decrease the density of compensating native defects via tuning the Fermi level position.36 We speculate that the hot-wall environment may play a role in an enhanced dissociation of molecular hydrogen to atomic hydrogen, which is beneficial for p-type doping. Details on exploring carrier gas composition for increasing the non-passivated Mg acceptor and free-hole concentrations are presented elsewhere.29 It is instructive to compare samples M3 and M6 with the same [Mg] but significantly different levels of non-passivated Mg of 1×1018 cm3 and 3×1018 cm3, respectively. The comparison reveals that M6 having three times larger fraction of non-passivated Mg exhibits four times higher free-hole concentration before annealing with regard to M3. Assuming ionization energy of 180–200 meV corresponding to NA of 1×1018 cm3 and 3×1018 according to Ref. 7, the maximum free-hole concentrations in the as-grown M3 and M6 layers are expected to be in the mid and high 1016 cm3 range, respectively. The p0 values measured for M3 and M6 are lower indicating a certain degree of compensation. We recall that M6 is grown at a higher V/III ratio as compared to M3, i.e., at lower supersaturation (Table I), and exhibits significantly lower levels of C (Table II and Fig. 3). Under the N-rich conditions, typically employed in MOCVD growth, CGa acting as a donor has the lowest formation energy in p-type GaN.12 Although [C] is at least two orders of magnitude lower than the non-passivated isolated Mg acceptors ([Mg]–[H]), it is close to the observed free-hole concentrations in the as-grown layers and will interfere negatively with the p-type conductivity. This can explain the observed lower free-hole concentration before annealing in M3 as compared to M6. After annealing, the free-hole concentration in M6 is also higher as compared with M3, correlating with NAND. In this case, NAND is dominated by the thermally activated Mg acceptors resulting from the dissociation of H from the Mg–H complexes in the as-grown samples.

The increase in Mg concentration above 2.4×1019 cm3 leads to a decrease in NAND, which becomes significantly lower than [Mg] (Fig. 7), consistent with previous reports.2,8 The C–V measurements show that 72%–92% of the incorporated Mg atoms are not electrically active. This doping range corresponds to [H] [Mg] in the as-grown samples as can be seen from Fig. 1 and Fig. S2 in the supplementary material and the respective fraction of Mg not bound to H is in the range of 74%–84% with no distinct dependence on [Mg] (Fig. 8). These results imply that the number of Mg atoms not binding to H in the as-grown layers could account to a large extent for the difference between the net acceptor and Mg concentrations in the respective samples after annealing. In other words, the fraction ([Mg]–[H])/[Mg] corresponds to scenario (ii) in which Mg incorporates at non-acceptor sites in the GaN crystal lattice.

This is reflected in a significant reduction of free-hole concentration in the annealed samples (Fig. 7). Note that the lower the NAND (Fig. 7) the lower the free-hole concentration in the as-grown (Table II) and annealed (Fig. 7) samples. Our STEM analysis shows that the observed large fraction of electrically inactive Mg could potentially be explained by the estimated concentration of Mg segregated at the PIDs. However, the fraction of Mg not bound to H is very similar for samples M5, M7, and M8 (Fig. 8) while their NAND differs significantly (Fig. 7). For example, sample M5 with an inactive Mg fraction of 84% has more than twice as high net acceptor concentration than M8 with 74% of inactive [Mg] non-bound to H, while the opposite could be expected. In fact, in M5, the concentration of Mg–H complexes before annealing is 9.6×1018 cm3 (assuming all available H binds to Mg), which compares well with the NAND of 1.7×1019 cm3 after annealing. This suggests that no significant concentrations of compensating donors, most notably VN and its complexes with Mg, are generated. In samples M7 and M8, the Mg–H levels before annealing are slightly higher as compared to M5, while NAND are significantly lower, i.e., (5–8)×1018 cm3, indicating that compensating defects are present in large densities. Both M7 and M8 are grown at a lower V/III ratio, which can explain an enhanced formation of VN in this case. Our findings strongly suggest that compensating donors, likely VN and its complexes with Mg are generated for highly Mg doped GaN grown at higher supersaturation and which play a significant role in the observed reduction of free-hole concentration (Fig. 7). The increase in [Mg] above 1×1020 does not alter the fraction of non-passivated electrically inactive Mg. In this case, however, the high [C] levels of 3.1×1017 cm3, due to the significantly lower growth temperature, lead to additional compensation and the GaN layer shows semi-insulating behavior.

It has been shown that by growing GaN:Mg on GaN substrates, with significantly lower density of dislocations, the free-hole concentration in the homoepitaxial layers can be significantly enhanced.38 To reduce the dislocation density in our heteroepitaxial case, we grew Mopt consisting of a 550-nm-thick GaN:Mg layer on 1-μm-thick undoped GaN buffer layer using the same growth conditions as for sample M3. Both screw- and edge-type dislocation densities of Mopt show the lowest values of 1.6×107 cm2 and 4.3×108 cm2, respectively, among all samples (Table II). NS and NE in GaN:Mg are likely to be even lower since the measured dislocation densities are averaged over the entire layer thickness. The free-hole concentration in the as-grown Mopt is increased more than twice as compared to the respective value in M3. If a similar approach is employed for lower Mg doping levels, e.g., such as in M2, free-hole concentrations in the low to mid 1017 cm3 range can be potentially achieved without the need of thermal activation. The free-hole concentration of 8.4×1017 cm3 and the resistivity of 0.77Ω cm in the annealed sample are among the best results reported in the literature.3,38 These results demonstrate that the hot-wall MOCVD capabilities can be expanded to deliver state-of-the art p-type GaN enabling development of power diodes and vertical transistors, and normally-off HEMTs, for example.

We have explored hot-wall MOCVD for the growth of Mg doped GaN. It is found that the efficiency of Mg dopant incorporation is similar to the earlier reported MOCVD results. A detailed study of growth and doping conditions’ impacts on the incorporation of Mg, H, and C is presented and the findings are discussed in terms of Ga supersaturation. Our results indicate that Ga supersaturation can be conveniently used as a universal parameter for the optimization of Mg incorporation and C reduction in MOCVD of GaN instead of multiple growth parameters independently used for the same purpose.

In the low doping range ([Mg] 2.4×1019 cm3), a large fraction (up to 21%) of active Mg acceptors (not passivated by H) is evidenced in as-grown material, resulting in free-hole concentrations in the low-to-high 1016 cm3 range. Hence, hot-wall MOCVD addresses a practical challenge related to the realization of as-grown p-GaN without the need for ex situ post-growth annealing for Mg activation. These results provide an intriguing opportunity to exploit the technique for delivering lightly p-type doped material required in device heterostructures, e.g., vertical MOSFETs,39 and simplifying the device fabrication process. In the annealed samples, NAND is found to closely follow [Mg], indicating that all Mg atoms are incorporated as acceptors, where 96% are activated by thermal dissociation of H during post-growth annealing.

In the high doping range above 2.4×1019 cm3, increased Mg leads to the generation of PIDs. It is shown that this can be correlated with the formation of hexagonal hillocks, the density and height of which increase with increasing [Mg]. This gradual surface morphology deterioration can be clearly observed in large-scale AFM imaging while on a smaller scale the morphology and surface roughness remain virtually unchanged. STEM analysis reveals segregation of Mg at the PIDs, sufficient to account for the observed amount of electrically inactive Mg not passivated by H. However, generation of VN and its complexes with MgGa need to be invoked to explain the considerably lower free-hole and net acceptor concentrations measured in the highly doped GaN:Mg grown at high supersaturation.

Under optimized growth conditions, high-quality GaN:Mg with a resistivity of 0.77 Ω cm and a free-hole concentration of 8.4×1017 cm3 corresponding to Mg acceptor activation of 5.25% has been demonstrated. These results and the established comprehensive picture of GaN:Mg growth via hot-wall MOCVD substantiate the expanded capabilities of the technique to deliver state-of-the-art p-type GaN for the development of power diodes and transistors.

See the supplementary material for additional details about Mg and H incorporation in the studied layers as well as on the structure of the observed pyramidal defects as revealed by STEM.

This work is performed within the framework of the competence center for III-Nitride technology, C3NiT—Janzén supported by the Swedish Governmental Agency for Innovation Systems (VINNOVA) under the Competence Center Program Grant No. 2016-05190, Linköping University, Chalmers University of Technology, Ericsson, Epiluvac, FMV, Gotmic, Hexagem, Hitachi Energy, On Semiconductor, Saab, SweGaN, and UMS. We further acknowledge support from the Swedish Research Council VR under Award No. 2016-00889, Swedish Foundation for Strategic Research under Grant Nos. RIF14-055, RIF14-0074, and EM16-0024, and the Swedish Government Strategic Research Area in Materials Science on Functional Materials at Linköping University, Faculty Grant SFO Mat LiU No. 2009-00971. The KAW Foundation is also acknowledged for support of the Linköping Electron Microscopy Laboratory.

The authors have no conflicts to disclose.

The data that support the findings of this study are available within the article and its supplementary material.

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