Change in Growth Mode of BGaN Layers Grown on GaN

Abstract

A change in the growth mode from Stranski–Krastanov one, which is characteristic of MOCVD grown GaN, to the laterally grown BGaN in the Volmer–Weber growth mode is described. This change in growth is evidenced by scanning electron microscopy (SEM) and transmission electron microscopy (TEM) images of BGaN grown on GaN at high temperatures. It is postulated on the basis of SIMS and XRD results that this change in growth is initiated by the transfer of boron atoms from gallium substitutional to interstitial. The proposed mechanism for the observed growth change is related to the generation of nitrogen interstitials and subsequent reactions with boron interstitials, which result in the formation of a BN layer at the growth front. The observed large change in the growth mode is due to a lattice mismatch between the grown BGaN and the atomic layer of BN and stays behind the change to the Volmer–Weber growth mode. The consequence of the Volmer–Weber growth mode is the textural layer of BGaN. The textural character of this material is associated with large voids between grown BGaN “plates”. These large voids are responsible for the termination of threading dislocations propagating in the c-direction. It is also postulated that the blocked threading dislocations from the GaN underlayer and laterally grown BGaN layers along the a-directions are responsible for the decrease in defect concentration within these layers.

1. Introduction

Ternary boron gallium nitride (BGaN) can be considered as a material to extend the bandgap engineering of nitrides. However, the main challenge to the epitaxial growth of BGaN is related to the low solubility of B in the GaN lattice. The upper solubility limit of B on a Ga sublattice (B_Ga) in GaN has been found to be only a few percent when grown by MOCVD and decreases with growth temperature, as was described previously [1]. The mechanism responsible for the decrease in B_Ga with growth temperature has been revealed by recent SIMS measurements showing that the total amount of boron is constant for all growth temperatures [1]. This suggests that an increase in growth temperature leads to the transfer of B atoms from the Ga substitutional sites to the interstitial positions. Thus, it was found that only the concentration of B_Ga depends on the growth temperature [1]. However, it was found that the growth mode of BGaN on GaN behaves differently from other nitrides.

It is known that there are three epitaxial growth modes: the so-called Frank–van der Merwe, Stranski–Krastanov, and Volmer–Weber [2]. The Frank–van der Merwe growth mechanism is observed in homoepitaxial growth or in heteroepitaxial growth with a very small lattice mismatch between the deposited material and the substrate. In heteroepitaxial growth with a larger lattice mismatch, layer-by-layer growth occurs in the initial stages, forming the wetting layer. However, as the deposited layer becomes thicker, the elastic energy forces the formation of three-dimensional islands or threading dislocations. Such formation of islands, or dislocations, reduces the elastic energy; this is the Stranski–Krastanov growth mode [3].

The Volmer–Weber mechanism is typical of highly mismatched systems between the substrate and the grown layer. Island growth occurs directly on the substrate without the formation of a wetting layer. The Volmer–Weber growth mode consists in the first phase of certain number of surface nuclei and in the second phase of their spreading.

These three mechanisms allow us to understand the growth of low-dimensional structures such as quantum wells and quantum dots. These growth modes are the result of the interaction between the substrate and the grown layer.

The epitaxial growth of GaN/Al₂O₃ or GaN/AlN is characteristic of Stranski–Krastonov one, leading to a density of threading dislocations in the GaN epitaxial layer of about 10⁹–10¹⁰cm⁻² [4]. Such a high concentration of threading dislocations is associated with the heteroepitaxial growth of GaN on AlN/saphire substrates. However, in the case of GaN growth on AlN single crystals, the dislocation concentration is much lower, about 10⁴cm⁻² [5].

Normally, ternary nitride layers on GaN, forming quantum wells, grow in the same growth mode as GaN one.

However, the growth of BGaN/GaN at high temperatures suggests a transition from the Stranski–Krastanov mode to the a-direction for BGaN in the Volmer–Weber growth mode. Such a transition has never been observed for nitride growth.

In the present work, we have discussed the mechanism responsible for such a change in the BGaN/GaN growth mode.

2. Experimental

GaN and BGaN layers were grown using the MOCVD method in a low-pressure metalorganic vapor phase epitaxy reactor, model AIX 200/4 RF-S as it was described previously in [1]. In the case of GaN, hydrogen (H₂) was utilized as the carrier gas, while trimethylgallium (TMGa) and ammonia (NH₃) acted as precursor gases. The flow rate of NH₃ was 2000 sccm and that of TMGa was 200 sccm for both GaN and BGaN. In the case of BGaN growth, the flow rate of TEB was 5 sccm. The same precursor gases and triethylboron (TEB) were used for the growth of BGaN. The epitaxial growth of GaN occurred at a temperature of 1090 °C and a pressure of 200 mbar. BGaN was grown at temperatures of 840 °C, 940 °C, 1040 °C, and 1090 °C, in identical pressures of the gas mixture. Growth of both materials used 2-inch single-sided polished Al₂O₃ wafers that were (0001)-oriented and covered with high-resistivity AlN. The reference sample in the study was 1.5 lm-thick GaN grown on AlN/sapphire at 1090 °C. In the case of BGaN, a GaN buffer layer was first deposited onto the AlN/sapphire template. The growth time of the GaN and BGaN layers under the above conditions was 20 min. Due to the same growth conditions, the thickness of the subsequent layers was 1400 nm and 600 nm for the GaN and BGaN layers, respectively. The determined roughness parameter (RMS) for the reference GaN layer was 1.45 nm. However, for BGaN layers, the surface roughness generally increases with the growth temperature and was 27.50 nm, 79.60 nm, and 61.7 nm for BGaN grown at 940 °C, 1040 °C, and 1090 °C, respectively [1]. Considering that the evaporation rates of different types of adatoms on the surface (B, Ga, and N) change differently with increasing temperatures, it can be assumed that the actual B/III/N adatom ratio on the surface is responsible for the increase in roughness of the growing BGaN.

Such grown BGaN layers have been investigated by several experimental techniques and results for these BGaN/GaN layers have already been published [1,6,7].

It was shown by X-ray Diffraction (XRD) that the amount of boron incorporated into gallium sites (B_Ga) depends on the growth temperature [1]. The XRD showed that B content changes from 2.5% to 0.73% for growth temperatures from 840 °C to 1090 °C, respectively. However, the total amount of boron within BGaN layers determined by SIMS was found to be constant for all growth temperatures [1]. That indicates that an increase in BGaN growth temperature leads to a transfer of B atoms from Ga substitutional sites to the interstitial positions. It was also shown that a significant accumulation of boron takes place at the BGaN/GaN interface [1].

It was also revealed, using various experimental techniques, that the growth temperature has an important influence on the electrical and optical properties of BGaN layers [6]. By using a high-resolution photoinduced transient spectroscopy (HRPITS) technique [8], several deep levels have been found within the BGaN layers [6]. The results of the HRPITS measurements also show that the increase in growth temperature has a significant influence on the decrease in defect concentrations. It was also found that improvements in mobility-lifetime product, increase in infrared luminescence, and increase in photoconductivity was found with an increase in growth temperature [6].

In publication [7], measurements of the photoinduced metastable EPR signal in BGaN layers were investigated. The intensity of the EPR signal grown at a temperature ranging from 840 °C to 1090 °C was compared with the results for trap energies and concentrations obtained in the HRPITS measurements for these samples. This comparison allows us to identify the complex responsible for the photogeneration of the paramagnetic state due to the vacancies complex [V_N-V_Ga]³⁺.

In the present work, the morphology of cross-section surface of BGaN layers is investigated using scanning electron microscopy (SEM) and transmission electron microscopy (TEM) measurements. The SEM investigation was performed using a Hitachi SU8230 cold field emission scanning electron microscopy. The imaging was conducted in deceleration mode (1 kV), with backscattered and second electron detectors. The TEM measurements have been performed using FEI Tecnai G2 F20 S-TWIN (FEI company, Hillsboro, OR, USA) operating at 200 kV on a cross-section sample prepared by the focused ion beam method.

3. Structural Properties of BGaN Layers

Despite the same growth conditions, the thickness of the subsequent layers was 1400 nm and 600 nm for the GaN and BGaN layers, respectively. The boron concentration in BGaN layers grown at 940 °C, 1040 °C, and 1090 °C, as determined by XRD, was 1.64%, 0.89%, and 0.73%, respectively [1]. These concentrations are associated with boron being on gallium substitution sites. On the other hand, the concentration of boron determined by SIMS in all samples was approximately 2.5% in all samples [1]. The discrepancy between these two results was explained as suggesting that a large fraction of boron escapes from gallium substitution sites with increasing growth temperature [1]. Using the HRPITS technique, six deep-level defects with activation energies of 570 meV, 640 meV, 780 meV, 900 meV, 1200 meV, and 1350 meV were found in the BGaN layers. The results of the HRPITS measurements also show that the increase in growth temperature has a significant effect on the decrease in defect concentration. The total defect concentration was found to be 1.7 × 10¹⁷ cm⁻³, 1.2 × 10¹⁶ cm⁻³, and 9.0 × 10¹⁵ cm⁻³ at growth temperatures 940 °C, 1040 °C, and 1090 °C, respectively [6].

The overview of the surfaces of GaN and BGaN samples by SEM has been shown in [1]. In contrast to the smooth surface of the GaN sample, the BGaN layers show a granular morphology, as discussed in our previous work [1]. XRD measurements also showed that these BGaN grains are well oriented along the c-axis. Therefore, they form the texture structure.

It is known that GaN heteroepitaxial layers on AlN substrates grow in the Stranski–Krastanov [3] growth mode. This growth mode is caused by a significant lattice mismatch between the layer and the substrate. The lattice mismatch between the substrate and the layer creates a built-in strain and, as a consequence of the increasing elastic energy, multiple threading dislocations are formed. The growth of the BGaN layer at 940 °C, as shown in Figure 1, although it may be perturbed by the strain introduced by the B-N bonds, which are considerably shorter than the Ga-N bonds, appears to follow the Stranski–Krastanov growth mode.

However, the more surprising results are obtained from SEM cross section images of the BGaN layers grown at higher temperatures, as shown in Figure 2 and Figure 3.

There is a clear change in the growth mode of the BGaN layer grown at 1040 °C, shown in Figure 2. The GaN layer clearly grows in the c-direction, which is typical for the growth of this material. However, above the GaN layer the BGaN tends to grow in the a-direction. Voids are formed between the layers of BGaN, perpendicular to the c-axis.

This change in growth mode is even more pronounced for the BGaN layer grown at 1090 °C shown in Figure 3. The grown BGaN layer is reminiscent of the Volmer–Weber growth mode. In a Volmer–Weber growth mode the observed growth consists of a certain number of surface nuclei in the first phase and their spreading in the second phase. Volmer–Weber growth results in a high degree of mosaicism of the material within the layer.

The BGaN layer shown in Figure 3 consists of “plates” perpendicular to the c-direction. This is a clear indication that the growth mode of the BGaN layer has changed from the Stranski–Krastanov growth mode to the Volmer–Weber growth mode. The characteristic for the Volmer–Weber growth mode is the observation of a certain number of surface nuclei in the second phase of their spreading in the a-direction. A high degree of mosaicism of the material within the BGaN layer is also observed. The BGaN “plates” along the a-direction are much better developed for the growth at 1090 °C than for the layer grown at 1040 °C.

The BGaN/GaN interface consists of a layer of approximately 150 nm, which appears to have been grown in the Stranski–Krastanov growth mode. This layer of BGaN is decorated with several micro voids, as can be seen in Figure 2 and Figure 3. The micro voids are only present in this interfacial layer, and the density of micro voids in this layer is in the range 10¹¹–10¹² cm⁻². Above this layer, there is a BGaN layer in which a drastic change to the Volmer–Weber growth mode takes place. In addition, the SIMS results show that the GaN/BGaN interface contains a strong peak of boron [1]. The relationship between the boron peak on the BGaN/GaN interface and the change to the Volmer–Weber growth mode will be presented in the Discussion section.

Figure 4 shows the results of the TEM studies of the cross-sectional sample grown at 1090 °C and oriented along the [1–100] GaN zone axis. The bright field scanning TEM is shown in Figure 4a. The image in Figure 4a shows the region containing sapphire, AlN buffer, the GaN layer, and BGaN “plates”. Consistent with the SEM results, the TEM results of the BGaN layer show the texture structure and mosaic character of this layer. The TEM image shows that threading dislocations (TDs) start at the AlN/sapphire interface. Most of the threading dislocations terminate at the AlN/GaN interface, while some propagate through the GaN layer. The imaging of the TDs using different diffraction conditions (Figure 4b,c) shows that the dislocations crossing the AlN/GaN interface have mostly 1/3<$11\overline{2}3$> Burgers vectors, making them mixed character TDs. Some of these dislocations terminate at the inner void surfaces in the BGaN layer (like the TD marked with the yellow arrow in Figure 4c), while others reach the large voids present in the Stranski–Krastanov/Volmer–Weber interface, leading to a reduction of the TD density inside the textured BGaN layer.

4. Discussion

The dramatic change in the growth mode of BGaN/GaN shown in Figure 3 and Figure 4 has never been observed in other nitride growth processes. The c-lattice constant of BGaN grown at 1090 °C is almost the same as the GaN one (there is a difference of about 0.1%), and therefore Stranski–Krastonov growth of BGaN on GaN is expected. In fact, this mode of growth occurs during the growth of approximately 150 nm BGaN. This thin BGaN layer consists of several micro voids. However, as shown in the SEM and TEM images shown in Figure 3 and Figure 4, the change from the Stranski–Krastanov growth mode to the lateral growth mode characteristic of Volmer–Weber growth takes place. This growth mode change occurs at a growth temperature of 1090 °C. This unique aspect of the growth change appears to be related to the role of boron in the GaN lattice, although that amount of boron is relatively small, in the order of 1–2%.

The role of boron in changing the growth mode of BGaN layers appears to involve the transfer of B_Ga from its gallium substitution sites to interstitial sites. This transfer may be induced by a significant strain generated by B_Ga within the GaN lattice. Since boron is a much smaller atom than Ga, the incorporation of B into the Ga site causes a symmetric inward displacement of the nearest neighboring N atoms. The distortion around the boron atom results in a B-N bond length of approximately 1.66 Å [9], which is significantly shorter than the Ga-N bond length of 1.95 Å. When the B_Ga atom is moved to an interstitial position, the local strain around it can be relieved, and a Ga vacancy is formed. At the high growth temperature, about 1% of the boron atoms move from the gallium substitutional position to the interstitial position, as shown by the SIMS results [1]. The formation of interstitial boron can follow the following reaction [7,9]:

B_Ga → B_i(+3) + V_Ga(−3)

(1)

The B_i(+3) ion consists of two electrons on a 1s shell and is therefore very small and can be very mobile. The mobility of B_i(+3) is confirmed by the SIMS results, showing the accumulation of boron atoms at the BGaN/GaN interface. This is in contrast to V_Ga(−3), which has a migration energy of 1.9 eV [10], indicating that it is not very mobile.

It may be expected that the other reactions take place in the vicinity of V_Ga(−3):

N_N → N_i(−3) + V_N(+3) or N_N → N_i(−) + V_N(+)

(2)

V_N has two stable charge states V_N(+1) and V_N(+3) [10,11]. These reactions have a relatively high formation energy in the bulk GaN material [9]. However, their formation energy may be reduced at defect sites. In particular, the formation of N_i(−3) and V_N(+3) may have lower activation energy near the V_Ga(−3) site. To release N_N in the vicinity of V_Ga(−3), three bonds need to be broken. In addition, the V_N(+1) charge state has a very high migration energy of 4.3 eV and the migration energy of V_N(+3) is 2.3 eV [9]. These high migration energies will ensure that this defect interaction will remain in some form of complex with the V_Ga(−3) due to Coulombic interaction. Coulombic attraction between these two types of vacancies can lead to the formation of vacancies:

V_N(+3) + V_Ga(−3) → void

(3)

This may be the origin of micro voids observed in SEM and TEM images in the 150 nm BGaN interface layer between GaN and textured BGaN. Accumulation of voids generated by reaction [3] may be caused by a local fluctuations of boron concentration. The most likely local boron fluctuations will create favorable sites for the generation of nitrogen interstitials and nitrogen vacancies, leading to an increase in the size of the micro voids. N_i generated by reactions [2] may enter the interior of micro voids and form strong bonds by N₂ molecules. The density of micro voids at the GaN/BGaN is close to 10¹² cm⁻². In addition, the TEM measurements shown in Figure 4b,c show that some threading dislocations may go inside of such micro voids and be terminated.

SIMS measurements indicate a high concentration of boron atoms at the interface [1]. On the other hand, SIMS did not detect boron atoms on the other side of the interface, inside the GaN layer. This suggests that there is efficient trapping of B_i(+3) at the interface. This trapping of boron ions may be associated with another reaction related to the generation of interstitial nitrogen (N_i) and the formation of BN. Coulombic attraction between B_i(+3) and N_i(−3) will lead to the reaction as follows:

B_i(+3) + N_i(−3) → BN

(4)

Diffusion of B_i(+3) to the growth front may play a crucial role in the change in growth from Stranski–Krastanov to the Volmer–Weber growth. Following the formation of micro voids within the BGaN layer grown in the Stranski–Krastanov mode, some of B_i atoms may also migrate to the surface of the grown layer. These can induce N_i generation according to the reaction [2] and resultant formation of BN layer on the growth front. The most likely BN layer will be 2D hexagonal form (h-BN). The bond length in h-BN is 1.47 Å, which is much shorter than the bond length in GaN which is 1.95 Å. Such a large lattice mismatch between the grown BGaN and the atomic layer of h-BN will not allow further wetting of the next grown BGaN layer. Therefore, the growth mode has to be changed to the Volmer–Weber mode. The formation of the BN layer in front of the growth explains why the change from Stranski–Krastanov to the Volmer–Weber growth only occurs only at high temperatures. The formation of BN requires the generation of N_i at the growth front and this process is thermally activated.

The other characteristic feature of the growth shown in Figure 3 is the textured nature of the grown BGaN layer, expressed by “plates” separated by large voids. The presence of such “plates” can be explained as follows: The average thickness of the BGaN “plate” is about 100 nm, which corresponds to 200 BGaN atomic layers. At high growth temperatures, about 1% of the boron leaves the gallium substitutional position and moves to the surfaces of the “plate”. The presence of boron on the surface of the “plates” is an ideal source for trapping of nitrogen interstitials and forming a BN monolayer. This allows the formation of an atomic layer of BN on both surfaces of the BGaN “plate”. The presence of the atomic layer of BN, possibly in a hexagonal form, on the surfaces of the BGaN “plates” may block the increase in the thickness of the BGaN layer along the c-axis. Continued growth of BGaN must originate elsewhere on the next “plate”, which will be the origin of the next BGaN “plate” grown along the a-axis. The presence of large voids between “plates” may be the main reason for the termination of the propagation of threading dislocations running in the c-directions within the BGaN material.

The absence of threading dislocations within the BGaN layer and the growth along the a-direction may be responsible for the significant improvement in the electrical properties of these layers grown at high temperatures. It is known that threading dislocations in GaN, resulting from the lattice mismatch between the substrate and the grown layer, are surrounded by a Cottrell atmosphere containing various types of lattice defects [4]. The textured BGaN layers, growing along the a-direction are not affected by the lattice mismatch between grown layer and substrate and therefore have no threading dislocations and no Cottrell atmosphere of defects. The concentration of defects in BGaN grown at high temperatures remains at the level of 10¹⁶ cm⁻³ [6]. Such a low concentration of defects is present despite the fact that the transfer of B_Ga to B_i is associated with the creation of about 10²⁰ cm⁻³ of gallium vacancies and boron interstitials [1]. However, “plates” of BGaN separated by large voids create regions where gallium vacancies can diffuse and be annihilated with nitrogen vacancies. On the other hand, a high concentration of B_i can be removed by the formation of BN layers.

The reduction of c-direction threading dislocations in GaN has been proposed previously. One method of reducing threading dislocations in GaN has been achieved by growth interruption followed by annealing in silane (SiH₄) [12]. TEM revealed the presence of pyramidal pits on the silane-exposed surface [12]. These pits were sites where dislocations drastically changed their direction of propagation from parallel to the c-axis to horizontal, resulting in a reduction in their density [12]. A technique to induce growth along the a-direction and reduce the dislocation density of heteroepitaxial GaN layers has been proposed using PENDO epitaxy [13]. The embedded voids approach (EVA), which relies on the generation of microvoids to reduce the dislocation density, has also been proposed [14]. These voids obtained by inductive coupled plasma-reactive ion etching provide free surfaces for dislocation termination [14]. All of these methods stimulate the growth of GaN along the a-direction, suggesting that the growth of nitrides along the a-directions is advantageous in terms of their properties.

5. Conclusions

The main result described in this paper is the experimental demonstration of the change in growth mode from Stranski–Krastanov one, the characteristic for MOCVD-grown GaN, to the Volmer–Weber growth mode of BGaN. This is observed in SEM and TEM images of BGaN/GaN grown at high temperatures. It is shown that this growth mode change is initiated by a large transfer of boron atoms in the BGaN layer from gallium sites to interstitials. This process stimulates the accumulation of boron interstitials at the growth front of the BGaN layer, and, at high temperatures, the generation of nitrogen interstitials and the formation of a thin BN layer. A large lattice mismatch between BN and BGaN will not allow the wetting of BN by the grown BGaN layer. This forces the change in growth mode from Stranski–Krastanov to the Volmer–Weber growth mode. As a result of the Volmer–Weber growth mode, the textural layer of BGaN is grown. The texture character of BGaN is associated with large voids between growth “plates”. These large voids are responsible for the termination of threading dislocations. Preventing dislocations from entering the textured BGaN layer stimulates the lateral growth mode mechanism and reduces the defect concentration in this layer.