Comparison of Chemical, Electronic, and Optical Properties of Mg

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Comparison of Chemical, Electronic, and Optical Properties of Mg-Doped AlGaN Jaakko Mäkelä, Marjukka Tuominen, Tiina Nieminen, Muhammad Yasir, Mikhail V. Kuzmin, Johnny Dahl, Marko P.J. Punkkinen, Pekka Laukkanen, Kalevi Kokko, Jacek Osiecki, Karina Schulte, Mika Lastusaari, Hannu Huhtinen, and Petriina Paturi J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b09248 • Publication Date (Web): 03 Dec 2016 Downloaded from http://pubs.acs.org on December 8, 2016

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The Journal of Physical Chemistry

Comparison of Chemical, Electronic, and Optical Properties of Mg-Doped AlGaN

Jaakko Mäkelä1, Marjukka Tuominen1, Tiina Nieminen1, Muhammad Yasir1, Mikhail Kuzmin1,2, Johnny Dahl1, Marko Punkkinen1, Pekka Laukkanen1*, Kalevi Kokko1, Jacek Osiecki3, Karina Schulte3 Mika Lastusaari4,5, Hannu Huhtinen6, and Petriina Paturi6

1

Department of Physics and Astronomy, University of Turku, FI-20014 Turku, Finland

2

Ioffe Physical-Technical Institute of the Russian Academy of Sciences, RU-194021 St.

Petersburg, Russian Federation 3

MAX IV laboratory, P. O. Box 118, Lund University, SE-221 00 Lund, Sweden

4

Department of Chemistry, University of Turku, FI-20014 Turku, Finland

5

Turku University Centre for Materials and Surfaces (MatSurf), Turku, Finland

6

Wihuri Physical Laboratory, Department of Physics and Astronomy, University of Turku, FI-

20014 Turku, Finland

ACS Paragon Plus Environment

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ABSTRACT: Hydrogen, carbon, and oxygen are common unintentional impurities of Al(x)Ga(1-x)N crystals. This impurity structure and its interplay with Mg impurities in Al(x)Ga(1-x)N semiconductors are relevant to develop the p-type nitride crystals for various devices (e.g, LEDs, transistors, gas sensors), but are still unclear. Here we have investigated Mg-doped Al0.5Ga0.5N before and after post-growth annealing with valence-band and core-level photoelectronspectroscopy, photoluminescence, and resistivity measurements. First, it is found that a surface part of the Al0.5Ga0.5N crystal is surprisingly inert with air and stable against air exposureinduced

changes.

Thus,

the

relatively

surface-sensitive

photoelectron-spectroscopy

measurements reflect in this case also the bulk-crystal characteristics. The measurements reveal the presence of deep states up to 1eV above valence-band maximum before and after the annealing, and that oxygen and carbon occupy N lattice sites (i.e., ON and CN). The model where CN -induced acceptor states in the band gap participate in the blue emission (photoluminescence) is supported. Furthermore, the presented Mg2p core-level spectra demonstrate that part of Mg atoms forms direct bond(s) with oxygen in the bulk-like structure of Al0.5Ga0.5N, and that the chemical environment of Mg atoms is much richer than was expected previously.

1. INTRODUCTION The Mg-doped Al(x)Ga(1-x)N semiconductor crystals have attracted great scientific and technological interest.1-16 Undoubtedly, the most important application of the p-type Mg-doped Al(x)Ga(1-x)N crystals has been the blue LEDs which have enabled the energy-efficient lighting.1-6 Furthermore, various applications of Mg-doped Al(x)Ga(1-x)N are under intensive research and


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development: for example, high electron-mobility transistors,17-21 ultraviolet LEDs,15,16,22,23 gas sensors,24-26 and electrodes of the photoelectrochemical cells.14,26,27

This progress has required surmounting of several challenges related to the synthesis and properties of III-V nitride materials. One of the main challenges has been how to create high enough concentrations of free holes to synthesize the p-type nitride crystals for devices. Magnesium is so far the only impurity that meets the criteria for commercial devices based on the nitrides. Still, an efficient Mg doping is a challenge, in particular, when the Al content of Al(x)Ga(1-x)N is increased.15,16 The Mg-doped crystals grown by common chemical vapor deposition (CVD) methods need to be post-growth treated in order to activate the hole doping. The as-grown Al(x)Ga(1-x)N materials include H impurities which passivate Mg-induced acceptor states, and therefore, the crystals are usually heated after the epitaxial growth, in order to remove/decrease the hydrogen passivation of Mg acceptor states. In addition to this established concept, there are also other plausible mechanisms which passivate or compensate the Mg doping because the achieved hole concentrations are still lower than what can be expected from the alloyed Mg concentrations.

Indeed there are several choices for compensating the p-type Mg doping. For example, nitrogen vacancies (VN) and oxygen substitutional sites (ON) can act as the compensating donor centers.8 Due to the CVD growth conditions, C and O are common impurities of Al(x)Ga(1-x)N materials in addition to the hydrogen. Because the Al-O bond formation is energetically favorable, the O concentration of the films can be expected to increase with increasing Al content. Both C and O can occupy three types of the crystal sites: the group-III lattice site; the N


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site; and/or any interstitial sites, depending on the Fermi-level position in the band gap and the ratio of group-III atoms to N in the growth conditions.28,29

A related open question concerns the origin of the blue luminescence (around 2.8 eV) which Mg-doped GaN usually emits, and which shifts toward the ultraviolet region when the Al content is increased (to about 4.7 eV for Mg-doped AlN).10,12,28-31 The blue luminescence (BL) has been associated with the electron transitions between donor-acceptor pairs (DAP), consisting of the ON or VN donor level and the Mg acceptor state. However, it has been recently12 found that also pure MgIII acceptors (i.e., Mg in the group-III lattice site) can cause BL via lattice relaxation and the transition: Mg0III + e- → Mg-III. In addition, the carbon substitutional sites (CN) have been shown to cause deep acceptor states,28 potential for the BL transitions.

Because the dimensions of Al(x)Ga(1-x)N crystals for applications decrease continuously in nanotechnology, it becomes more and more relevant to understand the impurity structure and local variation of Al(x)Ga(1-x)N crystals on atomic scale. To contribute to that target, we report here on the impurity structure of Mg-doped Al(x)Ga(1-x)N by combining core-level and valenceband photoelectron-spectroscopy, photoluminescence, and resistivity measurements of Mgdoped Al0.5Ga0.5N before and after post-growth annealing. The focus has been toward obtaining high-resolution signals from the atomic bonding environment of each element and electron states around the valence-band maximum using synchrotron-radiation photoelectron spectroscopy.


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2. EXPERIMENTAL SECTION Commercial Mg-doped Al0.5Ga0.5N film (500 nm thick), grown on a two-inch Si(111) wafer by CVD, was cut into rectangular pieces of 5 mm x 10 mm. The Mg concentration was higher than 1x1019 cm-3, but the as-grown film was not initially annealed above 500 °C in order to activate the p-type doping. Here it is worth noting that the as-grown sample was slightly degassed around 100-200 °C before photoelectron spectroscopy (PES) measurements to remove air-induced adsorbed species from the surface. This degassing temperature is however much lower than the hole activation temperature (>500 °C), and therefore the sample has been still labelled as-grown (or the sample before the annealing) in this work. After the first measurements, the hole doping of Al0.5Ga0.5N pieces was activated by annealing around 700 °C in vacuum conditions for 10 min, and then the samples were characterized again. For Mg2p core-level measurements, a reference sample was also prepared by depositing an approximately 3 nm thick Mg film on a cleaned Si(100) substrate. The Mg film was allowed to partially oxidize in air during the sample transfer to the synchrotron center. Photoelectron spectroscopy was performed at the MAX-lab in Lund on beamline I311. A sample holder with Ta clips was used for mounting the sample and the Fermi-level position was determined by measuring the Ta Fermi edge with the same parameters as the PES measurement on the sample. A separate x-ray photoelectron spectroscopy instrument (at University of Turku) was also used to estimate amounts of O and C contaminants of Mg-doped Al0.5Ga0.5N. Furthermore, because PES is a relatively surface-sensitive probe, it is necessary to realize surface effects such as air-induced contaminants and structural changes on the PES spectra, in order to make a justified interrelationship between the PES results and the PL and resistivity measurements of the film. To that effect, we characterized surface properties of


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the Mg-doped Al(0.5)Ga(0.5)N film also by low-energy electron diffraction (LEED) and scanning tunneling microscope (STM).

Core-level spectra were analyzed by decomposing each spectrum into emission components by means of the spectral fitting analysis. Each individual component in the Al2s, C1s, N1s, O1s spectra is assumed to be a single Voigt-profile peak, while each component in the Ga3d and Mg2p spectra consists of two Voigt peaks, i.e., it is a spin-orbit doublet. The Shirley background was removed from each spectrum before the fitting. The Lorentzian width of the Voigt peaks, which arises from a finite lifetime of a core hole, was 0.2-0.4 eV depending on the binding energy. The Gaussian-width part, including an instrumental resolution, non-homogeneity of the sample, and thermal broadening at room temperature, was allowed to vary. Spectral features, such as shoulders and asymmetries, were utilized as an argument to introduce additional components in the fitting. The minimum numbers of components, at which the spectrum can be reasonably reproduced, were used in the fittings. This approach led to the Gaussian width of 0.81.1 eV for Al2s, C1s, Mg2p, N1s, and O1s, as well as the Gaussian width of 0.6 eV for Ga3d. The spin-orbit splitting and branching ratio were 0.45 eV (0.30 eV) and 1.5 (2) for Ga3d (Mg2p), respectively.

The temperature-dependent resistivity measurements up to 800 °C were executed with a standard two-probe configuration connected to a Keithley 6487 picoammeter for a separate wafer piece. The copper wiring was made to the surface of the film with silver paste (distance between the connectors was 5 mm) and the samples were attached to an Inconel plate that was heated with a PVD Products infrared diode laser. Photoluminescence (PL) spectra were measured at


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room temperature using excitation at 266 nm (4.66 eV) from a TII Lotis Nd:YAG laser and detection with an Avantes Avaspec HS-TEC CCD spectrometer. The excitation beam was filtered with a band-pass filter (Newport 10LF10-266) and on the emission side the excitation was blocked using a high-pass filter (Standa S3S-25).

3. RESULTS AND DISCUSSION Figure 1a shows that the resistivity of the Mg-doped Al0.5Ga0.5N film decreases due to the post-growth annealing at 700 °C. The result is consistent with the concept that the H passivation of Mg acceptor states decreases with the annealing, thereby activating the p-type doping. Using a logarithmic resistance plot, we have determined the activation energy of Mg acceptor states to be 0.2-0.3 eV after the post-growth annealing, which is consistent with previous findings.7-11 The acceptor activation energy of 0.6 eV, estimated for the as-grown film, also agrees with previous results.7

PL spectra (Figure 1b) from the Mg-doped Al0.5Ga0.5N before and after the annealing are dominated by the BL-type emission around 320-340 nm, consistent with the previous results associated with donor-acceptor pair (DAP) transitions.7,10 Furthermore, the emission around 300 nm has been previously associated with the electron transfer from the conduction edge, or shallow donor site, to the Mg-acceptor level.7,10 It can be concluded that the PL emission does not change significantly due to the post-growth annealing, in contrast to the resistivity. This indicates that the activation of Mg acceptors is not related to BL, corroborating previous findings.7 The PL intensities in Figure 1b can be compared to each other before and after the annealing (a separate GaN sample was used to adjust the intensity scales). Lower excitation


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power (not shown) changes the PL spectra such that the intensity of the emission shoulder at 3.5 eV increases in relation to the 4.1 eV emission. Thus the transition channel behind the 3.5 eV luminescence is expected to become saturated sooner than the channel for 4.1 eV transition.

Figure 1. (a) Resistivity measurements of Mg-doped Al(0.5)Ga(0.5)N as a function of the measurement temperature for the as-grown film and after the post annealing. Resistance is plotted on a logarithmic scale as a function of 1000/T in order to estimate the activation energy of acceptor states: EA = 597 meV (1st anneal), EA = 300 meV (2nd anneal), and EA = 178 meV (3rd anneal). (b) Photoluminescence spectra from Mg-doped Al(0.5)Ga(0.5)N before and after the post-growth annealing.

Before the photoelectron analysis, a characterization of the Mg-doped Al0.5Ga0.5N film with LEED and STM is presented in Figure 2 because the surface properties are also relevant in understanding how relatively surface-sensitive PES results can be connected to the above bulklike properties of the resistivity and PL. First of all, it is unusual that LEED diffraction spots can be observed from the AlGaN sample after long air exposure (several months) without any surface treatment. Usually a semiconductor surface part which has been exposed to air becomes immediately oxidized and amorphous, and thus does not provide any LEED spot before a proper surface cleaning procedure. Therefore, the LEED observation (Figure 2b) means that an AlGaN surface is exceptionally inert with air and stable against oxidation-induced structural changes.


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Moreover, the LEED pattern shows that the AlGaN surface follows the (1x1) bulk-plane structure without any significant reconstruction.

Figure 2. (a) Large scale empty-state STM image from the Mg-doped Al(0.5)Ga(0.5)N film after the annealing around 700 °C in UHV; the voltage 4.56 V and current 0.05 nA. (b) LEED pattern from the Mg-doped Al(0.5)Ga(0.5)N film, which is similar before and after the annealing, with electron energy of 203 eV (c) Empty-state STM zoomed-in image from the same surface as (a).

The annealing of Mg-doped Al0.5Ga0.5N did not change the (1x1) LEED pattern significantly, but it simplified STM measurements. The STM imaging was not straightforward before the postgrowth annealing due to the insulating nature of the 4.5 eV energy-band gap film; the p-type doping was not yet activated by the annealing. In Figure 2a, STM images after the annealing around 700 °C in UHV are shown. The large-scale STM image shows a terrace-step structure, which is typical for an epitaxial film. In the zoomed-in image (Figure 2c), it is possible to see atomic ordering (area B) and also surface defects (area A) with white dots, arising from, e.g., AlOx.

As we will discuss in more detail below, the PES results show clear oxygen and carbon signals from Mg-doped Al0.5Ga0.5N before and after the post-growth annealing, although the AlGaN


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surface is inert with air. In fact, the O 1s intensity decreased only little (if nothing) as a function of different surface treatments (we tested the following surface treatments: post-annealing in UHV; and in NH3 flow; Ar-ion sputtering; and sulfur-based chemical etching). This supports that also the bulk part of Mg-doped Al0.5Ga0.5N includes oxygen atoms. This is, in turn, consistent with the appearance of LEED patterns (Figure 2) before any surface treatment. If only the surface was oxidized, as for the traditional semiconductor crystal exposed to air, we should not see any LEED spots from an amorphous surface, like AlOx and GaOx. Also the C 1s line is clearly present after the different surface treatments, supporting that a bulk part of Mg-doped Al0.5Ga0.5N includes both oxygen and carbon.

To recapitulate, the AlGaN surfaces provide an exceptionally good platform for investigating the bulk-like properties of the semiconductor using the non-destructive PES methods because (i) the air exposure does not cause the formation of an amorphous and strongly oxidized coating, and (ii) the surface atomic structure does not reconstruct as compared to the bulk-plane structure. Therefore, AlGaN surfaces appear to be an exception among semiconductor crystals. Clearly, the amounts of oxygen and carbon defects can be expected to increase somewhat toward the outermost surface of Al0.5Ga0.5N. Still, O and C atoms occupy the wurtzite lattice sites rather than form separate phases in the AlGaN film because the studied Mg-doped Al0.5Ga0.5N film is well crystalline, as can be deduced from Figure 2b. On the other hand, such surface enrichment of impurity atoms makes it possible to get a strong enough PES signal from defects which appear also in the bulk crystal of Mg-doped Al0.5Ga0.5N. The example PES spectra measured with different photon energies (i.e., different surface-sensitivity conditions) in Figure 3 support the above conclusion. Because these spectra do not change significantly when the surface sensitivity


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is changed, the dominant spectral components, which are analyzed and described below, arise from the bulk-like structure beneath the outermost surface layer(s).

Figure 3. Al 2s, N 1s, Mg 2p, and C 1s spectra from the as-grown film (i.e., before the annealing) as a function of the surface sensitivity of the measurement, which increases toward the top spectra.

The valence-band PES spectra from Mg-doped Al0.5Ga0.5N in Figure 4a reveal that the valence-band edge moves toward the Fermi level (0 eV) due to the post-growth annealing. That is, the Fermi level moves in the band gap about 0.6 eV toward the valence-band maximum (VBM), which is consistent with the above resistivity measurements showing that the p-type doping is activated, at least partially, during the 700 °C annealing. For the as-grown film, the Fermi level lies near the midgap position because the VBM is about 2 eV below the Fermi level and the band gap of Al0.5Ga0.5N is 4.5 eV. Both spectra in Figure 4a exhibit non-zero photoelectron signal between the Fermi level and VBM, indicating the presence of filled acceptor-type gap states.


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Figure

4.

(a)

Valence-band

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photoemissions

measured with 50-eV photons from Mg-doped Al(0.5)Ga(0.5)N before and after the annealing. (b) Valence-band spectra from Mg-doped Al(0.5)Ga(0.5)N after the annealing using two different photon energies.

Figure 4b shows the valence spectra from Mg-doped Al0.5Ga0.5N after the post-growth annealing with two different photon energies. The 700 eV photons provide more bulk-sensitive signal than the 100 eV photons. The photon-energy broadening is also larger at the 700 eV conditions, causing broadening of the emission features as compared to the 100-eV spectrum. This broadening is suggested to cause the emission tail above the Fermi level (up to -0.6 eV) in the 700-eV spectrum. The comparison of these spectra indicates that the acceptor band-gap states are not predominantly caused by surface structure or contamination, which is further consistent with the above surface characterization. Therefore, it can be concluded that there are band-gap states around 0.3 eV and 1.0 eV above the VBM in the bulk-like structure of Mg-doped Al0.5Ga0.5N. The former, shallow states agree with the acceptor activation energy determined above from the resistivity measurements, and can be associated with the Mg acceptors. In contrast, the states around 1 eV above VBM can be related to deep acceptors. Such deep levels appear already before the annealing (Figure 4a). This is consistent with the PL spectra (Figure


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1b) exhibiting the dominating DAP emission that does not change much due to the post-growth annealing. Thus, it can be suggested that the blue emission arises from the transitions via these deep acceptor states. The results below support that the deep states arises from CN impurities.

Figure 5. Al2s, Ga3d, and N1s spectra from Mg-doped Al(0.5)Ga(0.5)N before and after the post-growth annealing. The spectral

fittings

with

minimum

number

of

different

components are also shown. The photon energy was 1200 eV for Al2s and N1s, and 50 eV for Ga3d. For Ga3d, the peaks marked by solid black lines are the d5/2 peaks of the doublet components, while red dash lines represent the corresponding d3/2 peaks.

Figure 5 presents the Al2s, Ga3d, and N1s core-level spectra from Mg-doped Al0.5Ga0.5N before and after the post-growth annealing. The core-level spectra have been aligned according to the binding energies of the annealed film, requiring the same binding-energy shift of about 0.6 eV toward lower binding-energies for the as-grown samples, as obtained from the valence band spectra.


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The core-level spectra of the as-grown samples (Figure 5) include a small additional emission or tail at high binding-energy (BE) side, as compared to the spectra after the post-growth annealing. Because the high BE tails disappear after the annealing, they are associated with a surface-related phenomenon discussed in the end of this chapter. The spectra in Figure 5 show also that the bonding environments of Al and Ga atoms vary more than those for N atoms because the Al2s and Ga3d spectra include more emission components than the N1s spectra. This is the first indication that O and C impurity atoms are incorporated into the N lattice sites rather than the group III sites.

To clarify this issue more, we estimated the peak intensities with the traditional PES instrument using Mg Kα source (at University of Turku) and compared the intensity ratios, taking into account of the atomic sensitivity factors of different core levels. Because the atomic ratio of (Al+Ga)/(N+O) was higher than unity (1.1 - 1.3), the comparison supports that both O and C atoms occupy the N lattice sites for the stoichiometric AlGaN crystal. However, because the (Al+Ga)/(N+O+C) ratio is smaller than unity, part of C and/or O atoms can occupy also group III sites.

Thus, the different Al2s and Ga3d components are connected to three types of group-III crystal sites: that bonding with four N atoms (III-N); site with at least one C bond (III-C); and site with at least one O bond (III-O). The electronegativities of C (2.5), N (3.0), and O (3.5), are used to assign the different components in the Al2s and Ga3d spectra. We cannot resolve the number of the different O and/or C bonds in more detail, but the above categories are well justified because of clear energy separations (at least 0.5 eV) between the different components. It is worth noting


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that relatively high intensities of III-C and III-O components can arise from the surface enrichment behavior, as discussed above.

The N1s spectra in Figure 5 contain only one dominating component, which arises from the N atoms bonded to Ga and/or Al and/or Mg atoms. In addition to that, the N1s spectra include the second component at the high BE side, which is associated with N-C bonding type, because the previous calculations show that the CIII lattice sites are favored in some conditions.28 It is expected here that the presence of hydrogen broadens the different components rather than causes a separate component.

Figure 6. O1s, C1s, and Mg2p spectra from Mg-doped Al(0.5)Ga(0.5)N before and after the post annealing. The fittings with minimum number of different components are also shown. The photon energy was 830, 580, and 290 eV for O1s, C1s, and Mg2p respectively. For Mg2p comparison, a reference sample of MgO/Si was also measured. For Mg2p, the peaks marked by solid black lines are the p3/2 peaks of the doublet components, while red dash lines represent the corresponding p1/2 peaks.

The O1s spectra in Figure 6 show that oxygen bonding environment is surprisingly heterogeneous in Mg-doped Al0.5Ga0.5N. This can be however understood if O atoms occupy the


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N lattice sites (ON), and bond to different elements: Al, Ga, Mg, and/or C. The presence of hydrogen is again expected to contribute the broadening of components. In contrast, the C1s spectra (Figure 6) are more homogeneous than O1s. The main C1s component is associated with carbide-type bonding, consistent with the concept that carbon atoms mainly occupy N sites (i.e., CN).28 However, there is also the second component at the high BE side, which can arise from a CIII bonding structure. To estimate the total C concentration in the bulk, the Mg-doped Al0.5Ga0.5N film surface was strongly sputtered by Ar ions before XPS intensity comparison (at University of Turku). Comparing the C1s intensity to the N1s one, we can estimate that the C concentration in the bulk is higher than 1x1019 cm-3. To recapitulate, the above photoelectron results show that O and C impurities mainly occupy the N lattice sites.

The Mg2p spectra, plotted in Figure 6 alongside the Mg2p emission from a control sample of oxidized Mg/Si, demonstrate the formation of Mg-O bonds and that the direct bonding environment of Mg atoms varies surprisingly much (i.e., three different components with 1.5 eV overall separation) in the Al0.5Ga0.5N film before as well as after the post-growth annealing. The survey spectra in Figure 7, which have been measured with extremely surface-sensitive photon energy of 140 eV, show that the topmost Al0.5Ga0.5N surface part does not include Mg atoms before the annealing. This supports that the Mg2p spectra reflect the bonding structure in the bulk-like structure of Mg-doped Al0.5Ga0.5N.

Figure 7. Survey spectra including Ga3d, Mg2p, Al2p, and Ga3p emissions before and after the post-growth annealing. The photon energy was 140 eV.


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This variation in the chemical environment of Mg atoms in the bulk structure is consistent with the incorporation of ON and CN. It has been commonly considered that the nearest bonding environment of Mg consists of only N atoms (i.e., MgIII -N) combined with H atoms (e.g., H interstitials). However, due to the easy incorporation of O and C into Al(x)Ga(1-x)N during the crystal growth, the formation of MgIII -ON and MgIII -CN bonds can occur. Furthermore it can be expected that MgIII -ON and MgIII -CN bonds are durable during the annealing, as compared to the hydrogen passivation of Mg-induced acceptor states. Therefore, it is not surprising that MgIII -ON and MgIII -CN appear also after the post-growth annealing (Figure 6). Future calculations are desired to understand how the MgIII -ON and MgIII -CN bonds affect the energy level of Mginduced electron states.

A decrease of the hydrogen passivation of the Mg-acceptor states might be seen as an increase of the low BE emission of Mg2p in Figure 6 after the post-growth annealing, if the effective electronic potential around Mg atoms increases (i.e., becomes more negative) due to the hydrogen removal. Another process which might also lead to the increase of the Mg2p intensity at the low BE side is following. Because the concentration of CN and ON atoms decreases somewhat through the evaporation from the surface in vacuum conditions due to the post-growth annealing around 700 °C, it is possible that N vacancies (VN) form. The vacancies might further facilitate the diffusion of C, O, and/or N atoms from deeper bulk toward the surface part during the post-growth annealing. Concerning the Mg atoms, the loss of the nearest neighbor via the diffusion process can cause unsaturated (broken) bonds for Mg atoms. Such dangling bonds can


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cause an increase in the electron charge around Mg atoms, leading to the enhanced intensity at low BE side of Mg2p. Thus, it might be useful to investigate and develop further the post-growth annealing environment for Mg-doped Al(x)Ga(1-x)N films, in order to maximize the removal of the hydrogen passivation and simultaneously to minimize the formation of the compensating defects, ON and/or VN.

4. CONCLUSIONS High-resolution PES spectra support that significant amounts of O and C are incorporated into Al(x)Ga(1-x)N bulk crystal during the material growth, and indicate that both O and C mainly occupy the N lattice sites before and after the annealing. The valence-band spectra reveal the presence of acceptor-like states up to 1 eV above VBM. These states can be associated with the blue luminescence, which does not change significantly due to the annealing, in contrast to the annealing-enhanced p-type conductivity. The results support the model that the CN –induced states participate in the blue emission. Moreover, the presented Mg2p spectra demonstrate the formation of direct Mg-O in the bulk-like structure of Al(x)Ga(1-x)N before and after the annealing. Thus, the bonding environment of Mg atoms is much richer than was expected previously.

AUTHOR INFORMATION Corresponding Author * E-mail: [email protected]


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ACKNOWLEDGMENT We thank the MAX-lab staff for their assistance. This work has been supported by CIMO and Academy of Finland (project 259213), and the European Community’s Seventh Framework Programme CALIPSO.

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Insert Table of Contents Graphic and Synopsis Here:


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Comparison of Chemical, Electronic, and Optical Properties of Mg

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