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Modified pulse growth and misfit strain release of an AlN heteroepilayer with a Mg–Si

codoping pair by MOCVD

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2016 J. Phys. D: Appl. Phys. 49 115110

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1. Introduction

Group-III nitrides are promising semiconductor materials for a wide range of technological applications in optoelec-tronics such as high power, high frequency and high temper-ature electronic devices [1, 2]. The success in fabrication of InGaN-based blue light-emitting diodes (LEDs) pushed prog-ress towards another new generation of shorter wavelength optoelectronic devices, turning to AlGaN-based structures. The AlGaN-based ultraviolet (UV) LEDs [3], compared to traditional UV light sources, possess distinct advantages and broad applications in, e.g. water and air purification, steriliza-tion, medical treatment, white light illumination, and optical storage. To chase short wavelength requires higher Al content

in the AlGaN compound, which has the largest band gap in the nitride family (AlN ~ 6.2 eV). AlN is also endowed with remarkable properties such as high thermal conductivity, high chemical stability and high hardness [4]. Hence, the growth of high crystal quality AlN on foreign substrates like sapphire is necessary for providing a basal layer for the epitaxy of high Al content AlGaN. However, the difficulty of conducting hetero-epitaxy of high quality AlN directly leads to considerable epitaxial dislocation density in the AlGaN layer. Therefore, improving the technique for producing an AlN hetero-epilayer has been of wide concern.

As we know, the Al atom is strongly stuck to the surface and thus has weak surface diffusion, which results in a slow lateral growth rate of the AlN epitaxial layer [5]. It easily

Journal of Physics D: Applied Physics

Modified pulse growth and misfit strain release of an AlN heteroepilayer with a Mg–Si codoping pair by MOCVD

Abdul Majid Soomro1,2, Chenping Wu1, Na Lin3, Tongchang Zheng1, Huachun Wang1, Hangyang Chen1, Jinchai Li1, Shuping Li1, Duanjun Cai1 and Junyong Kang1

1 Fujian Key Laboratory of Semiconductor Materials and Applications, CI Center for OSED, School of Physics and Mechanical & Electrical Engineering, Xiamen University, Xiamen 361005, People’s Republic of China2 Institute of Physics, University of Sindh, Jamshoro, Sindh, Pakistan3 Xiamen Industrial & Commercial School, Jimei district, Xiamen 361024, People’s Republic of China

E-mail: dcai@xmu.edu.cn and jykang@xmu.edu.cn

Received 14 November 2015, revised 5 January 2016Accepted for publication 13 January 2016Published 17 February 2016

AbstractWe report the modified pulse growth method together with an alternating introduction of larger-radius impurity (Mg) for the quality improvement and misfit strain release of an AlN epitaxial layer by the metal–organic chemical vapour deposition (MOCVD) method. Various pulse growth methods were employed to control the migration of Al atoms on the substrate surface. The results showed that the pulse time and overlapping of V/III flux is closely related with the enhancement of the 2D and 3D growth mode. In order to reduce the misfit strain between AlN and sapphire, an impurity of larger atomic radius (e.g. Mg) was doped into the AlN lattice to minimize the rigidity of the AlN epilayer. It was found that the codoping of Mg–Si ultrathin layers could significantly minimize the residual strain as well as the density of threading dislocations.

Keywords: strain release, pulse growth, AlN, doping

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forms 3D islands instead of 2D layers, seriously affecting the crystal quality of the epitaxial layer and the AlN surface smoothness. The high temperature growth and pulsed flow method was proposed for enhancing Al migration on the sur-face, and has achieved better smoothness of the AlN epilayer [6]. However, the pre-reaction between the Al source (TMAl) and ammonia usually causes polycrystalline growth and unin-tentional incorporation of impurities, e.g. oxygen and carbon. Alternately switching on and off V-III precursors by pulsed flow was reported as an effective way to prevent this pre-reac-tion, but the growth rate of AlN was then reduced. Due to the unavailability of large size bulk AlN substrates, sapphire, SiC and silicon are currently foreign substrates frequently used for the deposition of the AlN epilayer. Significant misfit strain between AlN and the foreign substrate inevitably gives rise to damage of the perfect periodic crystalline lattice of AlN and hence results in a large number of defects such as threading dislocations. A low or middle temperature AlN buffer layer was introduced prior to high temperature AlN epitaxial growth for minimizing the misfit strain. But the temperature window for the buffer growth was found to be very narrow and the effect was limited. Since the emergence of defects such as dis-locations is an unwanted, deconstructive way to release misfit strain, the introduction of other types of defects like impurities is desired, that do not damage the lattice and can tolerate the strain by adsorbing the misfit energy by expanding the lattices.

In this work, modified pulse growth methods were sys-tematically studied for the quality improvement of the AlN epitaxial layer on sapphire by metal–organic chemical vapour deposition (MOCVD) and a technique for misfit strain release in the AlN epilayer was proposed by alternatively introducing larger-radius impurities. Various pulse growth methods were employed to control the migration of Al atoms on the sub-strate during the growth periods of the buffer layer and high-temperature epitaxial layers. The results reveal that the pulse duration and overlapping of V/III flux will enhance the 2D and 3D growth modes, respectively. In order to minimize the misfit strain, impurities of different atomic radius were doped into the lattice to minimize the rigidity of the AlN epilayer. It was found that the codoping of Mg–Si heterostructural layers could effectively minimize the residual strain and dislocation density.

2. Experimental details

The AlN epilayer was grown on sapphire (0 0 0 1) by MOCVD [7] with a Thomas Swan closely coupled showerhead (CCS) reactor, which is designed especially for reducing the unde-sired gas phase parasitic reaction. Trimethylaluminium (TMAl) and ammonia (NH3) were used as precursors for Al and N, respectively, and H2 was used as carrier gas. The optimized V/III ratio was obtained between 1500–2000 for the pulsed growth mode. Bis-cyclopentadienylmagnesium (Cp2Mg) and silane (SiH4) were used as sources of dopants. A low pressure of 50 torr was employed during the growth to enhance epilayer smoothness and crystal quality. Prior to AlN growth, thermal cleaning of the (0 0 0 1)-oriented sapphire

substrate was carried out under hydrogen ambient at 1050 °C for 10 min to remove native oxide from the surface. Then, the temperature was reduced back to 580 °C to carry out nitrida-tion of the sapphire surface. After nitridation, the temperature increased again to 1040 °C for a short time of the deposition of the nucleation layer (NL) with continuous growth mode, followed by the growth of a 5 nm–30 nm buffer layer. Finally, the growth of the 900 nm-thick AlN epilayer proceeded under higher temperature of about 1100 °C by using the pulsed growth method. When the Cp2Mg and SiH4 were induced alternately into the AlN epilayer, all the precursors (TMAl and NH3) were retained with the optimized V/III ratio under 1100 °C.

After the growth of samples, crystal quality was evaluated by x-ray diffraction (XRD) with a Panalytical X’pert PRO system. The surface morphology and surface roughness of samples were investigated by Normarski optical microscopy (Olympus BX51M), photography and atomic force micros-copy (AFM) using a Veeco Dimension 3100 system in non-contact scanning mode. Mg concentration in the sample was measured by using a Quad PHI 6600 secondary ion mass spectrometer (SIMS) system and the Cs+ ion beam was used as the primary ion source. Raman spectra were recorded using a Renishaw InVia Raman Microprobe equipped with a 532 nm laser. The laser spot was about 2 μm in diameter and the spec-trum was taken by averaging over 5 circles.

3. Results and discussions

3.1. Control of nitridation treatment of the sapphire substrate

Commonly used substrates for AlN epitaxial growth are sapphire, SiC or Si. Although SiC has less lattice mismatch (3.5%) compared to AlN, the high cost significantly restrains its promotion for commercial applications. In contrast, a Si substrate is plenty and has a lower price. However, its opaque feature at short wavelength light pulls it back from the appli-cation region of UV optoelectronic devices. Overall, sapphire (Al2O3) (a = 0.476 nm) still plays the dominant role as the for-eign substrate for AlN (a = 0.311 nm) heteroepitaxy, though a large lattice mismatch is present [8]. Similar to the epitaxial growth of GaN on sapphire, lattice rotation of the AlN epi-layer with respect to the sapphire lattice will take place by about 30° in order to minimize the misfit strain, as shown in figure 1; even though the residual strain is still considerable and various epitaxial techniques have to be introduced to min-imize the influence from this significant mismatch.

For the formation and epitaxy of AlN on sapphire, a trans-ition from Al2O3 to AlN should be accomplished to lower the adhesive energy of the upper AlN epilayer. Thus, the nitridation treatment of the sapphire surface becomes critical at the begin-ning stage prior to the entire AlN epitaxy. For this purpose, NH3 was used to nitridize the sapphire surface under a certain thermal condition and to form a wetting surface for AlN deposi-tion, as shown in figures 2(a) and (b). Similar to GaN growth, it was found that nitridation under high temperature (>1000 °C) will lead to a N-polar face growth of AlN [9]. Therefore, low

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temperature nitridation was performed in order to ensure the epitaxy of the Al-polar AlN. In our work, various nitridation times from 0 s to 410 s were studied for the optimization of the crystal quality of AlN.

XRD rocking curves obtained on symmetrical (0 0 0 2) and asymmetrical (10–12) planes were measured for sam-ples #A1–#A3 and the corresponding full widths at half maximum (FWHMs) are summarized in table 1. From these results, one can see that without nitridation (sample #A1), the FWHMs of both planes are larger than 2000 arcsec, indicating very poor crystal quality and misaligned grain orientation. When the nitridation treatment was introduced onto the sap-phire surface (#A2 for 250 s), obviously the value for (0 0 0 2) as well as for (10–12) is significantly reduced, almost half of that value without nitridation. In principle, the dislocation density is proportional to the FWHM of the (10–12) plane.

The nitridation treatment of the sapphire surface produces an important additional wetting layer for the adsorption of AlN molecules and their effective nucleation. In this way, the crys-talline quality becomes better and the orientation of seeds is well arranged along the [0 0 0 1] direction with less tilt. Meanwhile, the dislocation density decreases. As the dura-tion of nitridation increases up to 410 s (#A3), the FWHMs

Figure 1. Crystal structure and lattice alignment of the AlN epitaxial layer on the sapphire substrate. (a) Mismatch in the lattice of AlN and Al2O3, and (b) minimization of misfit strain with orientation rotation of the AlN lattice by 30°.

Figure 2. Schematics of growth processes of the AlN epilayer on sapphire. (a) Bare sapphire substrate, (b) nitridation of the surface of sapphire by NH3 under low temperature, (c) nucleation of AlN islands, (d) buffer layer under middle/high temperature with coalescence of nucleation islands, (e) 2D growth of the AlN epilayer under high temperature by Al migration enhancement using the MMEE method, and (f) epitaxy of the smooth AlN epilayer. Red lines stand for threading dislocations.

Table 1. XRD results of AlN samples under different nitridation times.

SampleTime of nitridation (s)

(0 0 0 2) (arcsec)

(10–12) (arcsec)

#A1 0 2049 2930

#A2 250 924 1551

#A3 410 380 936

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continuously decrease and the optimal nitridation condition is reached. It was also found that when the nitridation was carried out for extremely long times, an excessive N-rich con-dition will be present on the surface and the migration of Al atoms becomes worse. Hence, we keep to using this optimized nitridation condition and go on to investigate the following AlN nucleation and epitaxy.

3.2. Role of nucleation and the buffer layer

In principle, the simple nitridation of sapphire is not sufficient for absorbing the huge mismatch strain. Pre-growth of the buffer layer prior to epilayer deposition is then necessary for accommodating the misfit strain and providing lower forma-tion energy of the AlN adlayer. Here, a middle/high temper-ature (below or around 1000 °C) is used to initialize the nucleation of AlN seeds followed by the growth of the AlN buffer layer, using continuous mode (figures 2(c) and (d)). The role of the nucleation layer is to allow the formation of AlN seeds in uniform distribution and lower density on the nitridized sapphire surface, as illustrated in figure 2(c). This could also lead to the formation of large AlN islands in the early stage of growth, which may also result in less coales-cence stress. Usually, the temperature window for the effec-tive buffer layer is rather narrow. After a series of temperature

optimizations, it was found in our case that a temperature lower than 1000 °C could not well buffer the misfit strain, the reason for which will be discussed later. Thus, a middle/high temperature of ~1040 °C was found to be the critical window for the effective buffer layer. A set of samples with different buffer thicknesses, controlled by deposition times of 0 s, 10 s and 40 s, was prepared for comparison. Upon the buffer layer, a 900 nm AlN epitaxial layer was grown at 1100 °C and the same condition was used for all samples. From this, the evalu-ation of the influence by the buffer layer could be discussed.

XRD rocking curves of the as-grown samples were recorded on the symmetrical (0 0 0 2) and asymmetrical (10–12) planes and the corresponding FWHMs were measured, as shown in figure 3(d). The 0 s-buffer layer means that the AlN epilayer was directly grown in the absence of the buffer layer. One can see that the FWHM of (10–12) is about 1700 arcsec, much larger than those of other samples. This indicates that without the buffer layer, the direct epitaxial growth of the AlN layer on the simply nitridized sapphire is subjected to a greater mismatch strain. Consequently, this results in a large number of threading dislocations and the crystal quality appears poor. By a systematic study of the buffer thickness, it was found that the optimal time of the nucleation buffer layer is around 10 s (5–20 s). Under this optimized condition, the FWHM of the (10–12) peak decreases to about 800 arcsec, indicating a

Figure 3. AFM images (a)–(c) and FWHMs of XRD rocking curves (d) for samples at different growth times of the buffer layer: 0 s, 10 s and 40 s, respectively.

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significant improvement of crystal quality by minimizing the dislocation density. With the increase of the buffer time up to 40 s, the (10–12) peak could not be improved much. Instead, the (0 0 0 2) FWHM gradually increases. These trends dem-onstrate that the insertion of the buffer layer can reduce dis-location density in the AlN epilayer to some degree, but too thick a buffer will introduce additional screw dislocations. The 10 s-buffer could keep the nucleated AlN islands in appro-priate density and size (figure 2(c)) and effectively buffer the misfit strain from the sapphire substrate. This consequently leads to the 2D growth of upper AlN epilayers.

The surface morphology of the above samples was scanned by AFM and is shown in figures 3(a)–(c). One can see that for the 0 s sample, the surface is covered by tens of hexag-onal islands of hundred nanometers size and the rms is about 3.7 nm. The large size AlN islands reflect the low island density and the 3D growth mode. Due to the large distance between islands, the coalescence and flatness of the AlN is poor. In contrast, for the 10 s sample (figure 3(b)) the AlN surface becomes much smoother (rms = 2 nm) and no sur-face dislocation pits are observed. A more uniform nucleation density could lead to the smooth coalescence of islands and thus the AlN grains quickly turn to the 2D growth mode with smoother surface. However, when the buffer growth time goes up to 40 s (figure 3(c)), the roughness of the AlN epitaxial layer surface drastically increases (rms = 6.1 nm) and mean-while a large number of black dislocation pits appear on the topmost surface. This may be attributed to the high nucleation density and the short island-to-island distance, which results in significant stress during the lateral coalescence of islands. This type of additional compressive stress between the neigh-boring islands will give rise to the formation of dislocations. Moreover, the squash of islands will cause the tilt of islands away from the vertical direction, as illustrated in figure 2(d). As a result, the XRD (0 0 0 2) rocking curve exhibits a broader FWHM. Meanwhile, the bad alignment of islands due to the squash easily produces coalescing gaps which finally lead to

dislocations and a rough surface. Thus, to balance these two effects of misfit-strain release and squash stress, it is of impor-tance to insert an optimized buffer layer of appropriate thick-ness and growth temperature.

Raman measurements were further performed on the samples to investigate the residual stress of the AlN epilayer [10]. Theoretically, the AlN phonon peak shift could reflect the stress state of AlN qualitatively. It has been reported

Figure 4. Raman spectra of the AlN epilayer with a 40 s buffer (a) and the shift of the characteristic AlN-E2high peaks for samples at

different buffer times (b). This demonstrates the release of residual misfit stress by the buffer layer.

Figure 5. Diagrams of the pulse growth mode. (a) MEE method, (b) MMEE method where an overlap of V and III sources in each circle is used to modify the migration enhancement. (c)–(f) Pulse diagrams for samples #A4-7, respectively.

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that the E2high peak of bulk AlN under zero stress is at about

657.4 cm−1, as shown in figure 4. From figure 4(a), one can see that all the characteristic Raman resonance peaks of both sapphire substrate and AlN epilayer are completely visible and clear. This demonstrates the high crystalline quality of the as-grown AlN epilayer. From figure 4(b), it is found that the sample without the AlN buffer layer is subjected to a tensile stress from the sapphire substrate whereas the introduction of the buffer layer switches the stress state back to compressive. Beyond the critical thickness, the influence of the lattice mis-match on the residual stress becomes less important. Hence, the residual stress imposed on the epilayer mainly stems from the exogenous stress (mainly caused by the thermal expan-sion mismatch) and the intrinsic stress (mainly produced by the grain squash during island coalescence). Since the growth condition of these AlN samples was kept the same, the exogenous stress affected by the substrate should be sim-ilar. Therefore, the origin of the residual stress can be mainly attributed to the influence of the buffer layer. For small-size grains, the inherent tensile stress is inversely proportional to the grain size [11]. In the absence of the buffer layer, the initial size of an AlN grain is usually small and thus the epilayer is subjected to a large tensile stress. On the other hand, when the buffer layer is inserted, the grain size will increase with increasing temperature [12]. Under a middle/high temperature (>900 °C), the buffer will give rise to the formation of larger nucleation grains [13]. Consequently, the influence of tensile stress could be avoided and the epilayer is then subjected to compressive stress. However, a thick buffer layer will lead to grains that are too large and hence will increase compressive stress. As a result, the crystal quality of the AlN epilayer will then deteriorate. Furthermore, the forbidden phonon peak at 615 cm−1 was only observed for the sample with a 40 s buffer layer, as indicated by arrows in figure 4(a), which is known as an indication of the disorder of grain orientations [14, 15]. This fact again emphasizes the importance of the optimization of buffer thickness.

3.3. Pulse growth for migration enhancement

Due to the sticky feature of Al atoms, a higher growth temper-ature (>1300 °C) is regarded as the best solution to enhance the Al mobility on the surface and to achieve high quality AlN epilayers [5, 16]. Another way is to pulse the supply of V and III sources alternately to enhance the migration of Al atoms under middle/high temperature (<1200 °C) [6, 17–19], which is also called the migration-enhanced epitaxy (MEE) [20] method (figure 5(a)). The important physics of the tech-nique is based on the fact that the pulsed flow could effectively

prevent the pre-reaction of TMAl and ammonia and the sole supply of TMAl could enhance the migration of Al on the sur-face, which leads to smoothness of the AlN surface. However, the MEE method will cause some selective adsorption at the edge of terraces, which in turn restricts the improvement of surface smoothness. In order to overcome this shortcoming of MEE, the modified migration enhanced epitaxy (MMEE) [20] method was further proposed, which combines the advan-tages of the continuous mode together with the pulse mode, as shown in figure  5(b). The MMEE method could effec-tively avoid the occurrence of a pre-reaction and meanwhile enhance the migration of Al atoms by flexibly controlling the cycle period and the overlap of V and III source supplies.

On the basis of an optimized buffer layer, various pulse sequences (figures 5(c)–(f)) with MMEE scheme have been employed for AlN epitaxial growth under 1100 °C. The tech-nical details of pulse sequences and the FWHMs of repre-sentative XRD peaks of these samples are listed in table 2. The surface morphology of #A4–A7 samples was scanned by AFM and is shown in figure 6. For sample #A4 with the conventional MEE method, one can see that the FWHMs of (0 0 0 2) and (10–12) peaks are 170 and 1370 arcsec, respec-tively. Apparently, the high FWHM of the (10–12) peak indi-cates high dislocation density. Thus, the MMEE method was used to modify the growth condition by varying the ratio of pulse and continuous periods. In sample #A5, continuous growth of 2 s was introduced into the pulse period, as illus-trated in figure 5(d). The results show that the (10–12) FWHM effectively decreases to 1073 arcsec, whereas the (0 0 0 2) FWHM increases. From the AFM images (figure 6(b)), we found that the grain size and height become larger, strongly indicating the enhancement of 3D growth of AlN grains by the insertion of continuous growth. The increasing surface rough-ness (rms = 15.7 nm) further confirms this influence.

For comparison, the pulse period in sample #A6 was increased up to 3 s to balance the 3D growth and meanwhile enhance 2D growth, as shown in figure 5(e). It can be found that the (0 0 0 2) FWHM dramatically decreases whereas the (10–12) FWHM goes up. As illustrated in the AFM image (figure 6(c)), the grain size becomes larger whereas they appear together with some tiny grains in between. This means that the lateral migration of Al atoms is enhanced by prolonging the pulse period and this 2D growth leads to increasing grain size. On the other hand, the additional small grains between large ones are actually introduced during that 2 s of continuous period. These small grains will also resist the coalescence of large grains. As a result, the non-uniformity of AlN islands leads to even more rough surface (rms = 26 nm). From the trend of the above samples, we could further understand that

Table 2. XRD results of different pulse mode grown AlN samples.

Sample Method Time of period (s) (0 0 0 2) (arcsec) (10–12) (arcsec) Ds (cm−2) De (cm−2)

#A4 MEE 4 170 1370 6.3 × 107 2.2 × 1010

#A5 MMEE 6 256 1073 1.4 × 108 1.3 × 1010

#A6 MMEE 8 20 1533 8.7 × 105 2.8 × 1010

#A7 MMEE 5 59 844 7.5 × 106 8.5 × 109

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the long pulse period will enhance the lateral 2D growth and the grain size. However, too long a pulse period meanwhile leads to unevenness and higher inter-grain stress. In contrast, the continuous growth mode can enhance the 3D growth for larger grain height, while too long a continuous growth will introduce many new nucleation sites, which deteriorates the surface smoothness. Therefore, overlapping of the TMAl and NH3 pulses (continuous growth) is introduced to mini-mize the lateral inter-grain strains. Based on the above anal-ysis, one can conclude that only the moderate MMEE period with balanced ratio of pulse and continuous segments could effectively improve the AlN crystal quality as well as the surface smoothness. In light of this, sample #A7 was grown with adjusted parameters, where the pulse period is kept at 2 s and the continuous growth time is reduced to 1 s, as illus-trated in figure 5(f). We found that optimized crystal quality could be obtained (FWHM of (0 0 0 2) is 59 arcsec and that of (10–12) is 844 arcsec) and the surface smoothness of the as-grown sample is well improved (rms = 2 nm). The screw (Ds) and edge (De) dislocation densities of AlN epilayers could be estimated from XRD rocking curves as well as AFM images. As listed in table  2, one can see that sample #A7 has the lower dislocation density of only 8.5 × 109 cm−2 (For details, please see the supporting information (stacks.iop.org/JPhysD/49/115110/mmedia).). Furthermore, a transmittance measurement was employed to investigate the optical prop-erty, as shown in figure  7(a). In principle, the Fabry–Pérot oscillation in the transmittance spectrum could reflect the optical quality of the semiconductor thin film. One can see

that a sharp band edge absorption appears at about 210 nm, which directly corresponds to the optical band gap of AlN, indicating the high quality of the AlN crystal. The strong and uniform oscillation in the spectrum demonstrates the smooth interface and surface of the AlN epilayer.

3.4. Alternating codoping of Mg–Si impurities

By introducing impurities in III nitride semiconductor, one can obtain not only different conducting types, e.g. n- or p-type conducting layers, but also improvements of crystalline quality. Si and Mg are the most commonly used n-type and p-type doping elements in III nitrides. In the previous study of the doping issues on III nitride, it was found that the influence on crystal quality is sometimes positive and sometimes nega-tive, which mainly depends on how the impurity atoms are introduced into the lattice. For example, Si-doping in an AlN epilayer would introduce additional residual stresses or even cracks; and Mg-doping was employed to improve the crystal quality of laterally overgrown GaN by enhancing the lateral growth rate [21].

As aforementioned, the heteroepitaxial AlN layer on sap-phire is subjected to a compressive stress. The atomic radius of Si is 117 pm, very close to that of Al (118 pm). In contrast, the Mg atom has a larger radius (145 pm) [22]. In principle, the partial substitution of Mg for Al atoms in the AlN lattice could make the lattice larger for absorbing the residual compressive stress, as illustrated in figure 7(b). By using (C5H5)2Mg and SiH4 as Mg and Si sources, respectively, the AlN epilayer weas

Figure 6. AFM surface topographies of AlN of samples (a) #A4, (b) #A5, (c) #A6 and (d) #A7.

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doped. The doped AlN epilayer was grown on the undoped AlN basal layer (~950 nm) which was grown by the optimized MMEE method. In order to modulate the stress release, peri-odic heavy δ-Mg and δ-Si doping was simultaneously inserted upon the background Mg-doping, as shown in figure 7(c). A total growth of 20 cycles of Mg–Si δ-codoping was carried out in the upper AlN epilayer. The depth profile of Mg concentra-tion in the AlN epilayer was investigated by SIMS, as shown in the inset of figure 7(a). The concentration of Mg incorpora-tion in the AlN layer is about 9 × 1018 cm−3. Because the Si was δ-doped into the AlN epilayer (ultrathin), the concentra-tion could not be clearly detected. Electrical measurements such as I–V curves were taken for this sample and the results showed that without thermal annealing, Mg impurities in AlN could not be effectively activated. This means that the acti-vated p-type carriers from Mg dopants could be negligible.

Figure 7(d) shows a real-time interferogram of the AlN epitaxial growth with alternating Mg–Si-doping. One can see that after buffer layer growth, the basal AlN epilayer was grown with higher reflective intensity by the MMEE method. Seven oscillations of growth can be distinguished with stable

amplitudes, indicating smoothness and a thickness of ~950 nm. Then, the growth is switched to the stage of Mg–Si-doping. During this period, the reflected intensity remains high but rough, which is due to the heavy δ-doping. By calculating the oscillation width, the thickness of the doped-AlN epilayer is determined to be about 80 nm. This sample was scanned by an XRD rocking curve, as shown in figures 8(a) and (b). The FWHM of the (0 0 0 2) peak is about 246 arcsec and that of plane (10–12) is reduced to be about 610 arcsec. The signifi-cant improvement FWHM of (10–12) indicates the reduction of edge dislocation density by the introduction of δ-Mg–Si-doping. The insertion of periodic δ-doping actually breaks the continuity of the uniform AlN deposition and forms a local heterointerface, as indicated in figure  7(c). These interfaces could partially terminate the propagation of threading disloca-tions or force them to bend. As a result, the final dislocation density in the topmost surface could be largely minimized.

Measurement of the E2 Raman spectra peak was performed to investigate the stress state of the doped AlN epilayer, as shown in figure  8(c). Compared with the Raman shift in figure  4(b), it is found that the residual stress in the AlN

Figure 7. Alternating doping of Mg and Si for strain release in the AlN epilayer. (a) Transmittance of the AlN epilayer of sample #A7, showing the excellent optical property with high transparency beyond the ultraviolet band. The inset shows the SIMS profile of Mg concentration. (b) Atomic structure of substitution of Mg for Al in the AlN lattice, which could accommodate more strain due to the larger atomic radius of Mg than the Al atom. (c) Structure of δ-Si/Mg alternating doping in the AlN epilayer, and (d) reflectivity as a function of growth time of AlN epitaxy.

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epilayer has been significantly released and almost reaches a stress-free condition. This stress-release mechanism could be attributed to two critical effects: one is the partial substitution of Mg for Al in the AlN lattice, which expands the AlN:Mg lattice slightly to accommodate the mismatch strain from sap-phire; and the other is the additional polarization field gen-erated between the heavily doped ultrathin δ-Mg (holes) and δ-Si (electrons) layers, which could partially screen the spon-taneous piezoelectric field in AlN and hence loosen the rigid lattice. On the other hand, AFM scanning shows the improved smoothness of the surface of the Mg-doped AlN epitaxial layer, the rms of which is down to 0.47 nm (figure 8(d)). This could also be attributed to the introduction of Mg during the AlN epitaxy, which will prevent the pre-reaction of Al and N atoms before arriving at the surface. Thus, the lateral migra-tion of Al atoms on the surface is effectively enhanced and the surface smoothness improved. This alternating Mg–Si doping scheme for residual strain release could be applied in the p-type AlN conducting layer for better crystal quality or in the buffer layer beneath AlN epilayers for pre-release of strains.

4. Conclusions

In conclusion, a modified pulse growth technique was pro-posed for the quality improvement of AlN epitaxial layers by the MOCVD method and the alternating codoping with larger-radius impurities was introduced for misfit strain release. Various pulsed growth methods were employed to control the migration of Al atoms on the substrate during the growth periods of the buffer layer and high-temperature epi-taxy. The pulse growth mode and continuous growth mode were found to be closely related with the enhancement of 2D and 3D growth of the AlN layer, respectively. It was found that impurities of larger atomic radius could be doped into the AlN lattice to minimize the rigidity of the AlN epilayer. The alternating codoping of Mg–Si heterostructural layers significantly minimized the residual strain as well as the den-sity of threading dislocations. These approaches can also be transferred to heteroepitaxial growth of other semiconductors which are suffering from low surface migration and/or high misfit strains.

Figure 8. (a), (b) XRD rocking curves of symmetrical (0 0 0 2) and asymmetrical (10–12) peaks of the AlN:Mg–Si sample, respectively. (b) Raman spectrum of as-grown AlN:Mg–Si epilayer, which shows the release of misfit strain. (d) AFM scanning of the AlN surface, showing the improved surface smoothness with rms = 0.47 nm.

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Acknowledgments

This work was partly supported by the ‘973’ programs (2012CB619301, 2012CB619304 and 2011CB925600), the NNSF (61204101 and 61574116), the NSFF (2013J01025), the FRFCU (20720150027), and the ‘863’ program (2014AA032608) of China.

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