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Abstract
An AlGaN quantum well (QW) structure of a deep-ultraviolet (UV) light-emitting diode (LED) needs to be well designed for controlling its band structure such that the heavy-hole (HH) band edge becomes lower than the split-off (SO) band edge and hence the transverse-electric (TE) polarization dominates the emission for achieving a higher light extraction efficiency. Here, we report the discovery of un-intentionally formed high-Al AlGaN nano-layers right above and below such a QW and their effects on the QW for changing the relative energy levels of the HH and SO bands. The comparison between the results of simulation study and polarization-resolved photoluminescence measurement confirms that the high-Al layers (HALs) represent the key to the observation of the dominating TE-polarized emission. By applying a stress onto a sample along its c-axis to produce a tensile strain in the c-plane for counteracting the HAL effects in changing the band structure, we can further understand the effectiveness of the HALs. The formation of the HALs is attributed to the hydrogen back-etching of Ga atoms during the temperature transition from quantum barrier growth into QW growth and vice versa. The Al filling in the etched vacancies results in the formation of an HAL. This discovery brings us with a simple method for enhancing the favored TE-polarized emission in an AlGaN deep-UV QW LED.
© 2017 Optical Society of America
1. Introduction
The light source in the deep-ultraviolet (DUV) range (200-300 nm) can be fabricated with the ternary compounds of AlN and GaN [1, 2]. So far, the external quantum efficiencies of DUV light-emitting diodes (LEDs) based on AlGaN are typically lower than ~10% [3–5]. AlGaN has an anisotropic emission behavior. The crystal-field split energy of AlN is different from that of GaN, leading to different orders of energy levels among different valence bands. In GaN, the band edge of the heavy-hole (HH) band is the lowest, followed by the split-off (SO) band, and then the light-hole (LH) band such that its emission is dominated by the electron transition between the conduction band and the HH band, which results in the transverse-electric (TE) polarized emission, i.e., polarization perpendicular to the c-axis or along the surface of a c-axis grown sample. On the other hand, in AlN, the band edge of the SO band is the lowest, followed by the HH band, and then the LH band such that the emission is dominated by the electron transition between the conduction band and the SO band, which leads to the transverse-magnetic (TM) polarized emission, i.e., polarization parallel with the c-axis or perpendicular to the surface of a c-axis grown sample. Therefore, the dominating emission polarization of AlGaN depends on the Al content. Usually, a higher Al content leads to a stronger TM-polarized emission. Normally, the dominating emission polarization of an AlGaN layer changes from TE polarization into TM polarization when the emission wavelength is shorter than 300 nm. Because TM-polarized light propagates along the lateral dimension of a c-axis grown LED structure, its surface light extraction is less effective, when compared with TE-polarized light [6–10]. Therefore, efforts are needed for changing the TM-dominated emission in a DUV LED into TE-dominated emission for improving its light extraction efficiency.
It is noted that the structure of the valence bands of AlGaN is controlled not only by its Al content, but also by its strain condition [11–13]. Under a compressive strain, the band edge of the conduction band is elevated and those of the valence bands are also increased (reduced in the coordinate system for electron), leading to a larger band gap for blue-shifted emission. The increase range of a valence band edge is related to its effective hole mass. Because the hole mass in the SO band is smaller than that in the HH band, the increase range of the SO band edge is larger, when compared with the HH band. Therefore, when a strong enough compressive strain is applied to an AlGaN compound of a high Al content, the lowest hole energy state can be switched from the SO band edge into the HH band edge, leading to TE-dominated emission. Various methods for increasing the compressive strain in an AlGaN quantum well (QW) have been proposed. The technique of epitaxial lateral overgrowth of AlN can increase the compressive strain in the AlN buffer layer for transferring an extra compressive stress onto the subsequently grown layers, including the QWs for changing their valence band structures [14–17]. However, this technique requires a complicated procedure, including the lithography and regrowth processes. Also, the lateral overgrowth for coalescence normally requires a growth thickness of 10-15 µm, which makes the LED cost significantly higher. When an AlGaN DUV QW structure is grown on AlN substrate, the induced compressive strain onto the QW can be stronger, leading to the favored valence band structure for generating TE-dominated emission [18–21]. However, so far, the price of AlN substrate is still much higher than that of sapphire substrate. With the cost consideration for DUV LED fabrication, sapphire substrate is still a reasonable choice.
Changing the nanostructure of a QW is also proposed for controlling its band structure. A staggered QW structure (two stages of different Al contents in a QW) [22, 23] and an AlGaN/GaN/AlGaN QW structure [24] have been demonstrated for adjusting the subband coupling effect and hence changing the valence band structure. A simulation study proposed the modifications of the wavefunctions of electron and hole for the same purpose [25]. Various methods have been proposed for increasing the efficiency of a DUV LED, including the change of the p-type structure for increasing hole injection [26, 27] and surface modification for enhancing light extraction [28, 29]. Among them, controlling the band structure of a QW for generating TE-dominated emission is one of the most effective ways for enhancing the efficiency of a DUV LED. In this regard, changing the QW structure, including the composition variations of the well and barrier layers, is an effective method for controlling its band structure.
In this paper, we demonstrate a simple method for effectively changing the band structure of AlxGa1-xN/AlyGa1-yN (x < y) QWs grown on c-plane sapphire substrate with the emission wavelength shorter than 300 nm such that the TE polarization dominates the emission. The change of the QW band structure is caused by the unintentionally formed high-Al layers (HALs) on both sides of a QW. The HALs are formed through the processes of hydrogen back-etching during temperature ramping and the Al filling in the Ga-etched vacancies such that the Al content is increased. With the higher Al content in the HALs, when compared with that in a quantum barrier (QB), on both sides of a QW, the quantum confinement in the QW is increased for changing the valence band structure. Material analyses provide us with the evidences for the existence of the HALs. The comparison between the results of polarization-resolved photoluminescence (PL) measurement and simulation study confirms the change of the band structure for switching the dominating emission polarization. The PL measurement with a stress applied onto a sample along its c-axis brings us with an idea about the effectiveness of HALs in changing the band structure. In section 2 of this paper, the growth conditions of the three AlGaN QW samples are introduced. The material analysis results, including transmission electron microscopy (TEM) and X-ray diffraction (XRD), are illustrated in section 3. Then, the comparison between the results of PL measurement and simulation study is demonstrated in section 4. The PL measurement data and the corresponding simulation results when an external stress is applied onto a sample are shown in section 5. The discussions about the cause for the above observed phenomena are made in section 6. Finally, conclusions are drawn in section 7.
2. Sample growth conditions
Three AlxGa1-xN/AlyGa1-yN (x < y) QW structures are grown with metalorganic chemical vapor deposition (MOCVD) on c-plane sapphire substrate. The structures of those three samples are the same except that the growth conditions of three periods of QW are different. The epitaxial growth starts with a three-temperature deposition process for forming the AlN buffer layer. The low, medium, and high temperatures at 960, 1130, and 1235 °C, respectively, are used successively with the V/III ratios at 1500, 300, and 75 to form the AlN layers of ∼50 nm, ∼150 nm, and ∼2.5 μm, respectively. Then, an AlGaN template layer of ∼1 μm grown at 1170 °C with the V/III ratio at 1200 is deposited, followed by the three periods of QW. In sample A, the well layer is grown at 1000 °C for 90 sec. A QB layer is grown under the same growth conditions as those for the aforementioned AlGaN template layer at 1170 °C for 30 sec. It is noted that with TMGa flow for QB growth, its growth rate is much higher than that of QW growth, which uses TEGa as the source. The QW growth procedures of samples B and C are the same as that of sample A except the temperature and duration for growing a well layer. In samples B and C, a well layer is grown at 940 °C for 90 and 120 sec, respectively. On the top of the QWs in each sample, an AlGaN capping layer of ∼10-20 nm is deposited under the same growth conditions as those for the AlGaN template layer below the QWs. The pressure of the MOCVD growth chamber is kept at 60 mbar throughout the whole growth procedure. In Fig. 1, we show the temporal variations of growth temperature, NH3, TMAl/TEGa, TMAl/TMGa, and H2 flow rates during the growth of a QW period. For switching the growth condition for a QB into that for a QW or vice versa, we arrange a temperature ramping period of 270 sec and a subsequent nitridation period of 180 sec, during which H2 flow continues. The PL emission peak wavelengths measured from the top surface at room temperature are 272, 282, and 295 nm for samples A-C, respectively. The PL measurement is excited by a 213-nm UV laser from the top surface (the epitaxial side).
Fig. 1 Temporal variations of growth temperature, NH3, TMAl/TEGa, TMAl/TMGa, and H2 flow rates during the growth of a QW period.
3. Material analysis results
Figures 2(a)-2(c) show the high-resolution bright-field TEM images of samples A-C, respectively, around the QW portions. The two dark stripes in each TEM image correspond to two QW layers in a sample. Between the two QW layers, there is a thicker QB layer of light contrast. Above the upper QW, the light-contrast layer can be another QB or the capping layer. Below the lower QW, the light-contrast layer can be another QB or the template layer. Either right above or right below a QW, one can see an even lighter-contrast thin layer in each of the three samples. Because in a bright-field TEM image, a lighter contrast corresponds to a higher Al content, it is speculated that a QW is sandwiched by two thin HALs, which have even higher Al contents than those of the neighboring QB, capping, and template layers, in each of the three samples. Such HAL structures around the QWs can be more clearly seen in Fig. 2(d), in which a high angle annular dark-field (HAADF) image of sample A is shown. In an HAADF image, a darker contrast corresponds to a higher Al content. Therefore, the Al content distribution in the three-QW structure of each of the three samples can be schematically shown in the line-segment drawing right to Fig. 2(d). Such a multi-stage-QB QW structure can be further understood with the geometric phase analysis (GPA) image in Fig. 3(b) based on the TEM image in Fig. 3(a) of sample C. In Fig. 3(b), the essentially red rectangular region corresponds to the QW layer, which is compressively strained on the c-plane, as shown in Fig. 3(c). Figure 3(c) shows the line-scan strain distribution along the pink dashed line in Fig. 3(b). The essentially green regions at the top and bottom of Fig. 3(b) correspond to the QB layers, which are also compressively strained. Between the QW and either the top or bottom QB, an essentially blue thin layer of tensile strain can be clearly seen. The tensile strain distribution indicates that such a layer has a higher Al content, when compared with the QW and QB layers.
Fig. 2 (a)-(c): High-resolution bright-field TEM images of samples A-C, respectively, around the QWs. (d): HAADF image of sample A with the Al content distribution in the three-period QW structure being schematically demonstrated by the line-segment drawing on the right.
Fig. 3 (a): A TEM image of sample C around the QW portion. (b): GPA image corresponding to the TEM image in part (a). (c): Strain distribution along the pink dashed line in part (b).
To further understand the structure of such an HAL and then its effect on the emission behavior of a QW, we can use a multi-stage-QB QW structure for fitting the measured XRD data. Before such a fitting process, we need to know the crystalline parameters of the AlN buffer and AlGaN template layers as accurate as possible. These two layers are strained by the thick sapphire substrate. Figures 4(a)-4(c) show the reciprocal space mapping (RSM) images on the (105) plane of samples A-C, respectively. Here, in each image, the two elliptically-shaped red regions correspond to the features of the AlN buffer and AlGaN template layers. The lighter, greenish elliptically-shaped regions indicate the QW satellite peaks. The centers of all these features are vertically aligned, implying that the QWs are fully strained. From the locations of the central points (maximum intensities) of the AlN-buffer and AlGaN-template features in each part of Fig. 4, we can read the coordinates, Qx and Qz, and hence evaluate the lattice constants, a and c, for each sample through the relations of a = [4(h2 + hk + k2)/3]1/2/Qx and c = l/Qz with (h k l) = (1 0 5) [30]. The results are shown in Table 1. For fitting the XRD data with the multi-stage-QB QW structure, we also need to estimate the thicknesses of various layers in a sample as the reasonable initial guesses. Such estimations can be obtained from the TEM images shown in Fig. 2 and others (not shown) for other portions of the samples.
Table 1. Coordinates of the central points (maximum intensities) of the AlN-buffer and AlGaN-template features, Qx and Qz, in Fig. 4 and the corresponding lattice constants, a and c, of the three samples.
Figure 5 shows the XRD ω-2θ scan profile (Exp.) and its best fitting curve (Fit.) of sample A. The insert schematically shows the multi-stage-QB QW structure used for XRD fitting. For fitting feasibility, we assume that the Al contents in the AlGaN template, capping, and QB layers are the same in the same sample. Also, in each sample, the structures of the three periods of QW plus HALs are the same, including the thicknesses and Al contents. However, the structures of the upper and lower HALs of a QW can be different. The optimized XRD fitting results like that shown in Fig. 5 for sample A are also obtained for samples B and C (not shown). The layer thicknesses and Al contents used for the optimized fitting are shown in Table 2 for all the three samples. Here, one can see that the HAL thickness ranges from ~0.9 to ~1.2 nm. In each sample, the Al contents of all the HALs are close to 88%, which is significantly higher than that of the AlGaN template, capping, and QB layers (in the range of 73.1-74.4%). The Al content of the QWs in sample A is close to 50% while those in samples B and C are around 40% because of the higher growth temperature for sample A. The QW thickness of sample C is larger than that of sample B because of the longer growth duration in sample C. The successful XRD fitting with the multi-stage-QB QW structure further confirms the existence of the HALs.
Fig. 5 XRD ω-2θ scan profile (Exp.) and its best fitting curve (Fit.) of sample A. The insert schematically shows the multi-stage-QB QW structure used for XRD fitting.
Table 2. Thicknesses and Al contents of various layers in the three samples obtained from XRD fitting.
4. Optical characterization and simulation results
To understand the QW emission behaviors, particularly their polarization dependencies, we measure the edge-emission PL spectra with the setup shown in Fig. 6(a). The PL emission from a sample edge, which is excited by a 213-nm laser beam illuminated onto the surface of the sample near the edge, passes through a polarizer for separating the TE- and TM-polarized emissions. Figures 6(b)-6(d) show the PL spectra of the TE and TM polarizations of samples A-C, respectively. In each sample, the TE-polarized intensity is significantly higher than that of the TM-polarized intensity. The Fabry-Perot oscillations in those PL spectra make the reading of spectral peaks difficult. By using a Gaussian curve to fit a PL spectrum near its maximum level, we can obtain the PL spectral peak. The fitted spectral peaks of the TE- and TM-polarized emissions (λTE and λTM, respectively) and their difference (λTE - λTM) of samples A-C are listed in rows 2, 5, and 8, respectively, of Table 3. Here, one can see that in all samples, the TE-polarized peak wavelength is longer than that of the TM polarization. In other words, in each of the three samples under study, the band gap of the HH band is smaller than that of the SO band. This result is inconsistent with what usually expected for an AlGaN QW with the emission wavelength shorter than 300 nm. The smaller band gap of the HH band can be due to the change of the quantum confinement caused by the HALs.
Fig. 6 (a): Setup used for measuring the polarization-resolved edge-emission PL spectra. (b)-(d): PL spectra of the TE and TM polarizations of samples A-C, respectively.
Table 3. PL spectral peak wavelengths and simulated emission wavelengths of the TE- and TM-polarized emissions (λTE and λTM, respectively) and their differences of the three samples.
To confirm the effects of HALs on the band structures and hence emission behaviors, we perform the simulation study for comparing with the experimental data. In the simulation study, one-dimensional 6 × 6 k‧p Schrodinger equations are solved for evaluating the band structures and emission characteristics of the three samples under study [11, 31–33]. By using the three-period, multi-stage-QB QW structure shown in the insert of Fig. 5 and the parameters shown in Table 2, we can evaluate the energy subband levels of the QWs, including the effects of the band structure modifications due to the strain and quantum confinement. From the QW subband structures, we can obtain the wavelengths of the TE- and TM-polarized emissions, which correspond to the transitions related to the HH and SO bands, respectively. The simulation results are shown in rows 3, 6, and 9 of Table 3 for samples A-C, respectively. Here, one can see that all the simulated emission wavelengths are quite close to the corresponding PL measurement results. In particular, the TE-polarized wavelength is longer than that of the TM polarization in each sample. For comparison, we remove the effects of the HALs by replacing the HAL Al contents by that of the QB in the regions of the HALs and perform the simulation. The results of the TE- and TM-polarized emission wavelengths in this situation are shown in rows 4, 7, and 10 of Table 3. Without the effects of HALs, the TE-polarized emission wavelength becomes shorter than that of the TM-polarized emission in each sample. In other words, without the HALs, the SO band edge becomes lower than the HH band edge and hence the corresponding band gap is smaller. When the HALs are added to the system, the quantum confinement condition of the QWs is changed such that the band gap of the HH band becomes smaller, when compared with the SO band. In this situation, the emission wavelength of the TE-polarization becomes longer.
5. Stress level to balance the effects of the HALs
Increasing the compressive strain in an AlGaN QW by controlling the structures and compositions of the underlying substrate, buffer, and template layers is a widely used technique for enhancing the favored TE-polarized emission. The formation of the HALs in the current study can also lead to the enhancement of TE-polarized emission. It is instructive to understand the effectiveness of the HALs in changing the band structure by estimating the level of an applied stress onto the QWs for balancing the HAL effects. For this purpose, we measure the polarization-resolved, edge-emission PL spectra when a pressure is applied onto a sample along its c-axis, as schematically demonstrated in Fig. 7(a). The pressure onto the sample is produced by a one-axis translation stage. Figures 7(b)-7(d) show the PL spectral peak wavelengths of the TE- and TM-polarized emissions under different pressures of samples A-C, respectively. The stress values in the abscissa are obtained by replacing a sample by a pressure gauge in the translation stage. In each sample, at zero applied pressure, the TE-polarized emission wavelength is longer than that of the TM-polarized emission, as shown in Table 3. As the applied pressure increases, both the TE- and TM-polarized emission wavelengths are elongated in each sample. However, the increasing slope of the TM-polarized emission is relatively larger such that the difference between the two polarizations becomes smaller as the pressure increases. In sample A (B), the curves of the two polarizations intercept at the stress level of 4.1 (9.8) kPa. In sample C, up to the maximum applied stress of 10.6 kPa, the TE-polarized emission wavelength is still slightly longer than that of the TM-polarized emission. The applied pressure onto the sample along the c-axis produces a compressive strain along the c-axis or a tensile strain in the c-plane onto a QW. In other words, the tensile strain caused by the applied stress can counteract against the band structure change produced by the HALs such that the band gap of the SO band can become smaller for emitting longer-wavelength light. In rows 2, 6, and 10 of Table 4, we repeat the emission wavelengths of the TE- and TM-polarized emissions at zero applied pressure. Then, in rows 3, 7, and 11 of this table, we show the corresponding values when the applied stress is 10.6 kPa. If we use the simulation results in rows 4, 7, and 10 of Table 3 to represent the emission behaviors under the condition without the effects of the HALs, we can see that up to 10.6 kPa, the applied stress has not completely balanced the effects produced by the HALs. We need an applied stress higher than 10.6 kPa for the complete balance. In other words, the HALs provide us with a quite strong effect to change the band structure of such an AlGaN QW for switching the emission dominance from the TM-polarization into TE-polarization. In this situation, the light extraction efficiency of an LED based on such a QW structure can become higher.
Fig. 7 (a): Setup used for measuring the polarization-resolved edge-emission PL spectra under a pressure along the c-axis of a sample, which is produced by a one-axis translation stage. (b)-(d): PL spectral peak wavelengths of the TE- and TM-polarized emissions under different stresses of samples A-C, respectively.
Table 4. PL spectral peak wavelengths of the TE- and TM-polarized emissions (λTE and λTM, respectively) and their differences of the three samples under zero and 10.6 kPa stresses.
6. Discussions – mechanism for forming the high-Al layers
As shown in Fig. 1, when the growth temperature ramps from 1170 down to 1000 or 940 °C for QW growth, the flow of H2 carrier gas continues although its flow rate is gradually reduced from 3800 to 200 sccm. During this process, hydrogen can back-etch the grown AlGaN layer. Because the Ga-N bonding energy of 2.2 eV is smaller than that of the Al-N bonding energy (2.88 eV), the back-etching effect mainly breaks the Ga-N bonds and relieves Ga atoms while preserving Al atoms, as schematically shown in Figs. 8(a) and 8(b) [34–36]. As shown in Fig. 1, during the following nitridation process, hydrogen back-etching continues. Meanwhile, nitrogen atoms incorporate, as demonstrated in Fig. 8(c). Then, when the Al and Ga atoms are supplied for QW growth, Al atoms not only can contribute to form the new AlGaN layer, but also can fill in the vacancies left by etched Ga atoms. Because of their shorter migration lengths [37, 38], Al atoms can more effectively fill in the vacancies, when compared with Ga atoms. In this situation, an AlGaN layer of a higher Al content is formed right below a QW. Such a process can also occur when the QW growth is completed and the growth temperature is ramped up for QB growth. The hydrogen back-etching effect during the ramping and nitridation processes can produce vacancies by relieving Ga atoms. Again, when the vacancies are mainly filled by Al atoms when Al source is supplied for growing a QB, an AlGaN layer of a higher Al content is formed right above a QW. Therefore, the key to forming an HAL is the hydrogen back-etching effect during growth temperature ramping.
Fig. 8 Schematic demonstrations of the atomic structure in the process for forming an HAL. (a): Before hydrogen back-etching. (b): After hydrogen back-etching. (c): After hydrogen back-etching and nitridation. (d): After the resupplies of Al and Ga sources.
Based on the XRD fitting results, the Al contents of all the six HALs in each of the three samples are about the same (~88%), as shown in Table 2, even though the QW growth conditions of the three samples are different. In particular, the Al contents of the HALs on the two sides of a QW are about the same even though those of the back-etched layers are different. Although the XRD fitting process cannot guarantee the unique solution, almost the same Al content in all the HALs of the three samples provides us with certain reasonable speculations. The Al content of an HAL can be controlled by the growth conditions, including the supply rates and durations of the metalorganic sources, NH3, and H2. If the duration of hydrogen back-etching is long enough, a maximum Al content in an HAL can be reached under a growth condition. The Al content of ~88% can be the achievable maximum level under the current growth condition. The formed HAL thickness is only 2-3 monolayers. This can be the maximum thickness of a layer, in which the lattice structure still exists even though many Ga atoms are removed such that Al atoms can fill in for forming an HAL [see Figs. 8(b)-8(d)]. Deeper etching can lead to the complete destruction of a superficial layer, which cannot contribute to an HAL. In other words, the process of hydrogen back-etching for forming an HAL is only suitable for the growth of a thin HAL or a superlattice structure. In LED application, the HALs on both sides of a QW may produce potential barriers hindering the electron and hole captures by the QW. However, because the HALs are so thin, the tunneling effect should lead to effective QW capture of electron and hole.
7. Conclusions
In summary, to understand the cause for the unexpected result of dominating TE-polarized emission in a DUV AlGaN QW structure, we first identified the existence of the higher-Al AlGaN thin layers right above and below a QW through TEM and GPA observations. Then, with XRD measurement and fitting, we obtained the thicknesses and Al contents of various layers in such a multi-stage-QB QW structure. Next, we compared the TE- and TM-polarized emission wavelengths or the corresponding HH- and SO-band related band gaps between the results of PL measurement and simulation to confirm that the existence of the HALs was the key to changing the dominant emission from TM-polarization into TE-polarization. By applying a stress onto a sample along its c-axis for producing a tensile strain in the c-plane, we could estimate the equivalent strain level in the QWs for balancing the HAL effects. The formation of the HALs was attributed to the hydrogen back-etching of Ga atoms during the transition of the growth temperature from QB (QW) into QW (QB) growth and then the vacancy filling of Al atoms when the Al supply was resumed. This discovery brings us with a simple technique to change the band structure in an AlGaN DUV QW for enhancing favored TE-polarized emission.
Disclosures
The authors declare that there are no conflicts of interest related to this article.
Funding
Ministry of Science and Technology, Taiwan (grants of MOST 104-2622-E-002-031-CC2, MOST 105-2221-E-002-159-MY3, MOST 106-2221-E-002-163-MY3, and MOST 105-2221-E-002-118); Excellent Research Project of National Taiwan University (NTU-ERP-105R89095A); US Air Force Office of Scientific Research (grant AOARD-14-4105).
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