Background

InGaN/GaN-based high-brightness light-emitting diodes (LEDs) and laser diodes, as the representative devices of III-nitrides, have attracted much attention owing to their important role in digital signage, high-density optical storage, and general illumination [1,2,3,4,5,6,7,8,9,10]. Generally speaking, fabrication of blue or green LEDs requires relatively high indium composition of InGaN layer [11, 12]. Although the reduction of growth temperature and the increase of growth rate of the quantum well (QW) can alleviate indium atom desorption to obtain high indium content, these methods also deteriorate the optical performance of InGaN/GaN multiple quantum wells (MQWs) by worsening interface abruptness and introducing more defects [13, 14]. Moreover, these defects usually act as nonradiative recombination centers, thus weakening the internal quantum efficiency of the device [15,16,17,18,19]. Therefore, achieving required indium content while maintaining high material quality is still a big challenge.

In order to settle the problems mentioned above, various growth techniques have been employed in striving for smooth morphology and sharp interfaces within the InGaN/GaN stack. Quantum barriers (QBs) grown at elevated temperature [20, 21] and growth interruption after QWs [12, 22] are widely used to improve the morphology of InGaN/GaN heterostructures. However, they all have their own limitations. For instance, barriers grown at high temperature may lead to severe In loss [14, 23]. Although growth interruption can improve morphology as well as reduce inclusions, it is at the expense of the optical quality of the QWs [21]. Recently, it is reported that introducing a small amount of hydrogen during the growth of GaN barriers can improve both optical and interface properties [24,25,26,27,28]. However, the effect mechanism of H2 on surface evolution of InGaN/GaN MQWs has not been fully understood yet.

In this paper, the effects of H2 proportion, defined as H2 flow divided by total carrier gas flow, during GaN barrier deposition, on surface morphology evolution are systematically investigated. Ga surface-diffusion rate, stress relaxation, and H2 etching effect are suggested to be the three main factors, affecting surface evolution. The dominant factors and their influences on the surface evolution are comprehensively discussed, which provides a technical guideline to obtain high-quality InGaN/GaN heterostructures.

Methods

The InGaN/GaN MQW structures were grown on c-plane sapphire substrate by Aixtron TS300 metal organic chemical vapor deposition system. Trimethylgallium (TMG), triethylgallium (TEG), trimethylindium (TMI), and ammonia (NH3) were used as precursors. Silane (SiH4) was used as the n-type dopant source. The structure was composed of 3.2-μm-thick undoped GaN layer and nominally six-period 2.4-nm-thick InGaN QWs separated by 11-nm-thick lightly Si-doped (n-doping = 3×1017cm−3) GaN barriers. A 1.0-nm-thick low temperature GaN cap layer (LT-GaN) was deposited immediately after the growth of QW layer. InGaN wells and GaN barriers were grown at 730 and 850 °C, respectively. A conventional InGaN/GaN MQWs sample, labeled as S1, was grown in nitrogen atmosphere. Four other samples, denoted as S2, S3, S4, and S5, were grown with different proportion of H2 flow to total carrier gas (N2 + H2) during barriers deposition, with the other growth parameters the same with S1. The percentage of H2 was 2.5% (S2), 6.25% (S3), 10% (S4), and 50% (S5), respectively.

The structures of InGaN/GaN MQWs were characterized by PANalytical Empyrean high resolution x-ray diffraction (HRXRD) system. Surface morphology was obtained by atomic force microscopy (AFM) (SPA-300HV) using tapping model. Room temperature (RT) photoluminescence (PL) properties of the samples were studied by 226-nm Nd-YAG laser with an excitation power density of 1.36 W/cm2.

Results and Discussion

The HRXRD ω-2θ scanning results of S1–S5 are illustrated in Fig. 1a. The strongest peak located at the center belongs to the underlying GaN template, and the satellite peaks correspond to the periodicity of the MQWs. It is found that the full-width at half-maximum (FWHM) of the strongest peaks in all samples is almost the same, indicating the similar crystal quality of GaN buffer layers for all samples. The presence of clearly distinguished “ + 4th” diffraction peak in samples S2–S4 manifest the improvement of crystal quality under low H2 percentage. The appearance of the “ + 5th” diffraction peak (represented by the rectangle in Fig. 1a) and the minimum FWHM value of InGaN “−1st” diffraction peak indicate the best interface quality in sample S3. The structure parameters determined by fitting the measured XRD curves are summarized in Table 1. The period thicknesses of the five samples are almost the same, and the values keep around 14.4 nm. The indium contents of the InGaN wells for samples S1 to S4 keep around 11.8%, while the value drops to 9.9% for S5. A large amount of H2 may etch the GaN LT-cap layer and then react with indium atoms in QWs, which result in the reduction of average indium content [29]. The roughness of the interface can be calculated by fitting FWHM of the XRD satellite peak by the following equation [26, 30]:

Fig. 1
figure 1

a The HRXRD ω-2θ scanning results of S1–S5. b The FWHM as a function of the satellite peak order and its linear fitting for the five samples

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Table 1 Structure parameters of InGaN/GaN MQWs determined by HRXRD fitting
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$$ \varDelta {\omega}_{\mathrm{n}}=\varDelta {\omega}_0+\left[{\left( \ln 2\right)}^{1/2}\cdot \kern0.3em \varDelta {\theta}_{\mathrm{M}}\cdot \frac{\gamma}{D}\right]\cdot n $$
(1)

where Δω n represents the FWHM of the n-th satellite peak, Δω 0 is the intrinsic width of satellite peaks, Δθ M is the angle spacing between the adjacent satellite peaks, D is the period thickness of the InGaN/GaN MQW and γ is the interface roughness. Figure 1b shows the linear relationship between FWHMs and satellite peak orders. The slope of the fitting line is related to the QW/QB interface roughness. The fitting results show that interface roughness is gradually reduced as the H2 percentage increases, and the optimum value is achieved at 6.25% of H2 (S3), as shown in Table 1. With further raising in the percentage to 50% (S5), the interface roughness is increased dramatically. Hence, the ratio of H2 during barrier growth has great impact on interface quality. A small percentage (0–6.25%) of H2 is favorable to obtaining sharp interface, while a large amount of H2 (50%) seriously roughens the interface.

The AFM images of sample S1–S5 are shown in Fig. 2a–e. The dark points are mainly V-pits [14, 31], which initiate at the threading dislocations (TDs) [21, 27]. The root mean square (RMS) surface roughness under different H2 percentage is illustrated in Fig. 3. The reference sample S1 grown with H2-free condition possesses the coarsest surface with an RMS roughness of 1.028 nm. The RMS value decreases with the increase of H2 percentage, and achieves the minimum value (0.705 nm) at 6.25% of H2, as shown in Fig. 3. As the H2 percentage raises to 10%, the surface gets slightly rougher. With further increase in H2 percentage to 50%, many large holes are formed, as pointed out by red arrows in Fig. 2, and surface RMS roughness dramatically increases to 0.924 nm.

Fig. 2
figure 2

The AFM images (10 × 10 μm) of five samples: a S1, b S2, c S3, d S4, and e S5

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Fig. 3
figure 3

The variation trend of pits density and RMS surface roughness under different H2 percentage in carrier gas during the growth of barriers

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Figure 4 shows the statistical calculated diagram of pits size distributions for the five samples. It can be seen that as 2.5% H2 is introduced, the smallest pits (<60 nm) start to emerge, and the largest pits (>160 nm) disappear. As H2 percentage increases to 6.25% (S3), the proportion of pits at size 80–100 nm is significantly raised, and that of large pits (>140 nm) is dramatically reduced to the minimum value. With further increase in the H2 percentage to 10%, the largest pits begin to emerge again. When 50% H2 is introduced, the ratio of large pits is dramatically increased. Hence, the pits size can be reduced by introducing a small amount of H2, and the optimum value is acquired at 6.25% percentage. However, the pits size shows an increase trend as H2 percentage further rises.

Fig. 4
figure 4

The distribution of pit size for the five samples: a S1, b S2, c S3, d S4, and e S5

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It is obvious that the evolution trend of RMS surface roughness is highly consistent with that of pits size, which may relate to the growth mode affected by the formed pits. Once the pits are formed, indium atoms will first nucleate at the point where the TDs intersect the InGaN/GaN interface [32,33,34,35], then island growth starts and finally island growth mode transfers to 2-Dimensional growth. In other words, the presence of V-pits will delay the 2-dimensional growth, then roughen the surface. The larger the size is, the more obvious the delay can be.

In order to elucidate the surface evolution mechanism under different H2 percentages, the variation trend of V-pits size is discussed in detail. As H2 percentage increases from 0 to 6.25%, the decrease in V-pits size possibly comes from the following two parts. First, the formed Ga-H complex may enhance the incorporation efficiency of Ga atoms on \( \left\{10\overline{1}1\right\} \) plane [35]. It is reported that the adsorption energy of the Ga-H complex is about 1.2 eV smaller than that of single Ga adatoms [28]. Hence, the attachment of hydrogen to Ga adatom could significantly weaken the bond to the surface, which benefits the surface diffusion of Ga atoms [28, 36]. Another important reason is the gradually enhanced etching effect with the increase of H2 percentage. Shiojiri et al. reported that indium atoms can be easily trapped and segregated around the core of TDs, which plays a role of a small mask that hinders Ga atoms migration [37]. Hence, introducing H2 during the growth can effectively eliminate indium-rich clusters at InGaN/GaN interface, and contribute to surface migration of Ga atoms [37,38,39]. In addition, hydrogen can etch some unstable areas, such as dislocation sites and V-pits [40,41,42,43]. It is reported that dislocation sites are unstable due to the high strain energy and weak binding energy, and these sites can be easily dissociated during the etching process [41,42,43]. Moreover, the V-pit commonly consists of six symmetric N-terminated \( \left\{10\overline{1}1\right\} \) facets [44, 45], which is much weaker during the etching process as compared with Ga-terminated facets [42, 43]. Therefore, when H2 arrives at the surface, it is difficult to etch most of the GaN on the surface due to the high stability of Ga-face. Thus, H2 etching occurs mainly at dislocation sites and V-pits [42, 43], causing the decomposition of GaN. Due to the low growth temperature of GaN barrier, the decomposition effect of GaN is weak when hydrogen percentage is low [26]. Hence, the enhanced Ga atoms incorporation plays a dominate role in surface evolution, which is beneficial in reducing the size and density of pits and in turn enhances the 2-Dimensional growth and suppresses the formation of new pits, and finally conductive to smooth surface. The correlation between pits density and H2 percentage is presented in Fig. 3. It is shown that the highest pit density (1.69 × 108 cm−2) exists in the H2-free sample. While a small amount of H2 is added in the carrier gas, pits density is gradually reduced and reaches the lowest value (0.92 × 108 cm−2) in sample S3. With further increase in H2 proportion to 50%, pits density is significantly increased to 1.28 × 108 cm−2. These results indicate that adding a little H2 in the growth of barrier layer can suppress the formation of new pits. However, the suppression of new pits formation could lead to strain accumulation inside the layer, and the strain may relax via formation of new dislocations and other defects such as big pits in S4 and S5 [21], which will deteriorate the quality of the surface, as well as the InGaN/GaN interface.

It is worth to mention that the large holes (>200 nm) as marked with red arrows do not appear in samples S1 and S2, and they only start to appear as H2 percentage becomes larger than 2.5%. The hole size in S5 is much larger than that in samples S3 and S4, which may relate to the following two possible mechanisms about hydrogen over-etching mechanisms. One is the hydrogen over-etching on dislocation sites and V-pits. As aforementioned, the enhanced diffusion of Ga atoms plays a dominant role when hydrogen percentage is low. However, this leading role shifts to the enhanced GaN decomposition around dislocation sites and V-pits when large amounts of hydrogen are applied. The hydrogen can diffuse along the dislocation line and then etch the surrounded unstable sites both vertically and longitudinally, which could decrease the average indium contents in MQWs region, degrade well/barrier interface quality and also form large holes on the surface. Another possible mechanism is about hydrogen over-etching on LT-GaN cap layer. As H2 proportion lower than 2.5%, the H2 etch effect on LT-GaN cap layer is negligible. As the H2 percentage increases to 10%, the H2 etch effect on LT-GaN cap is illustrated in Fig. 5a. H2 only etches a part of the cap layer, which has little influence on the QW layer, as evidenced by almost unchanged indium contents, and positive influence on the surface morphology, as confirmed by low pits density and small size of the holes. However, under large H2 percentage, the LT-GaN cap layer may be partly etched away and QW layer be directly exposed to H2, as presented in Fig. 5b. Under this case, H2 will react with indium atoms in QW layer, leading to significant indium loss, large size, and high-density holes, and consequently dramatically deteriorates InGaN/GaN interface and surface qualities. Hence, surface morphology evolution is an integrated effect of surface diffusion rate, strain relaxation and H2 etch effect. For sample S2 and S3 with H2 percentage lower than 6.25%, gradually enhanced surface diffusion rate and H2 etch effect play a dominant role, which contributes to smoother surface and lower pits density. With further increase in the percentage to 10% (S4), surface properties become slightly worse as a result of stress relaxation. The surface morphology of sample S5 grown in 50% H2 is mainly controlled by H2 over-etching effect and strain relaxation in InGaN QWs, which leads to many large holes and the worst surface.

Fig. 5
figure 5

a The etch effect on LT-GaN cap layer with H2 percentage lower than 10%. b the H2 over etch effect on LT-GaN capping layer under large H2 percentage

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Figure 6a shows the measured room temperature PL spectra of the five samples. It can be seen that the PL intensity shows an increase trend and peak energy exhibits blue-shift as H2 percentage increases. Compared with that of sample S1 without H2, the integrated PL intensity of samples S2–S5 is increased by 7.0, 15.8, 19.3, and 31.6%, respectively. For samples S2–S4, slightly blue-shifted peak energy and reduced FWHM are observed, as shown in Fig. 6b. As aforementioned, the structure parameters of sample S2–S4 are quite similar. Hence, the slightly changed spectral characteristics along with enhanced PL intensity are mainly caused by the enhanced surface and interface quality, and the partial relaxation of stress in QWs alleviating quantum confined stark effect (QCSE) [21, 46]. In contrast, the significantly reduced FWHM, blue-shifted peak energy and enhanced PL intensity of sample S5 may result from strain relaxation and the lowest indium content caused by H2 over-etching effect, both of which can greatly alleviate QCSE effect in MQWs [46,47,48,49]. In addition, H2 can eliminate impurities such as carbon and oxygen in active region, which would benefit the improvement of the PL intensity [50, 51].

Fig. 6
figure 6

a The measured room temperature PL spectra of the samples. b Gaussian-fitted FWHM and peak energy versus H2 percentage in carrier gas

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Conclusions

In summary, the effect of H2 percentage during the barriers growth on InGaN/GaN MQWs properties has been systematically studied. As a small percentage of H2 (≤6.25%) is introduced, the combined effect of enhanced H2 etch effect and surface diffusion contribute to the improvement of surface, interface and optical properties. In spite of the strongest PL intensity achieved by introducing large percentage H2 (50%), the integrated effect of H2 over-etching and stress relaxation degrades surface and interface quality of the InGaN/GaN MQWs. Hence, the use of H2 with appropriate proportion during the barriers growth can achieve smooth surface with low pits density and enhanced optical performance. The profound discussions of surface evolution mechanism here clearly depict the physical pictures of surface evolution process under different growth conditions, which is helpful for the fabrication of high-quality GaN-based devices.