Impact of Mg content on native point defects in Mg x Zn 1 − x O ( 0 ≤ x ≤ 0 . 56 )

Impact of Mg content on native point defects in MgxZn1−xO (0 ≤ x ≤ 0.56) J. Perkins,1 G. M. Foster,1 M. Myer,2 S. Mehra,2 J. M. Chauveau,3,4 A. Hierro,5 A. Redondo-Cubero,6 W. Windl,7 and L. J. Brillson1,8,9,a 1Department of Physics, The Ohio State University, 191 West Woodruff Ave., Columbus, Ohio 43210, USA 2Columbus School for Girls, 56 S. Columbia Ave., Columbus, Ohio 43209, USA 3Centre de Recherche sur l’Hetero-Epitaxie et ses Applications, Centre National de la Recherche Scientifique (CRHEA-CNRS), Rue B. Gregory, F-06560 Valbonne Sophia Antipolis, France 4University of Nice Sophia Antipolis, Parc Valrose, F-06102 Nice Cedex 2, France 5Dpto. Ingeniería Electrónica and ISOM, Universidad Politécnica de Madrid, Ciudad Universitaria s/n, 28040 Madrid, Spain 6Dpto. Física Aplicada y Centro de Micro-Análisis de Materiales, Universidad Autónoma de Madrid, 28049 Madrid, Spain 7Department of Materials Science and Engineering, The Ohio State University, 2041 College Road N., Columbus, Ohio 43210, USA 8Department of Electrical and Computer Engineering, The Ohio State University, 2015 Neil Avenue, Columbus, Ohio 43210-1272, USA 9Center for Materials Research, The Ohio State University, Columbus, Ohio 43210, USA

2][3] The 3.37 eV band gap of ZnO can be tuned by incorporating either Mg or Cd 4 to enable complex heterostructures that can enhance transport properties in ZnO based transistors 5,6 and optoelectronic efficiency of UV lasers, 7 light emitting diodes, and solar blind detectors. 8The ZnO band gap increases with Mg alloying and Mg can be incorporated into ZnO at low concentrations without significantly disrupting the wurtzite structure. 70][11] However, ZnO occurs naturally in a zincblende or wurtzite structure, while MgO occurs naturally in a rocksalt structure, and little study has been directed at the detection and characterization of defects in the (Mg,Zn)O alloys, especially with high quality, isostructural single crystals across a wide alloy range.
The two most thermodynamically stable defects in ZnO are oxygen vacancies (V O ) and zinc vacancies (V Zn ), particularly under our Zn-rich growth conditions. 12(Details of our growth conditions and defect thermodynamic stability are reported in Ref. 13.).5][16][17] Thus, V Zn defects act as acceptors to partially compensate degenerate carrier densities in Ga-doped ZnO. 18Similarly, defect complexes associated with V O can be associated with increased n-type doping in ZnO. 19These vacancies appear to be mobile since depth-resolved cathodoluminescence spectroscopy (DRCLS) reveals pronounced segregation of V O 20,21 and cation vacancies V C , either V Zn or V Mg (this paper).Besides doping, defect levels deep within the band gap of ZnO and (Mg,Zn)O represent traps that increase "non-radiative" recombination of free carriers for transport and optoelectronic applications as well as introduce interface states that affect Schottky barrier heights.Indeed, while (Mg,Zn)O alloys are envisioned as lattice-matched confinement layers for ZnO quantum well lasers, deep level defects would degrade such lasing as well as alter the heterojunction band offsets that determine quantum well depths.Here, we describe the spatial distributions, densities, and energy levels of V C and V O defects and their dependence on Mg alloy composition.These defects produce energy levels deep within the (Mg,Zn)O band gap that produce nearly all midgap luminescence intensity, reflecting their recombination velocities relative to band-to-band recombination and, in turn, the impact of these native point defects on transport and light emission efficiency.
We performed these studies using a wide (0 ≤ x ≤ 0.56) range of single crystal Mg x Zn 1−x O alloy compositions grown on r-plane sapphire by molecular beam epitaxy (MBE).This wide range is particularly useful in identifying electronic trends that would not be evident over a more limited range.Furthermore, X-ray diffraction (XRD) shows these a-face (Mg,Zn)O alloys are isostructural and single-phase over this range, 22 notwithstanding the Mg-rich lattice transformation from wurtzite to cubic rocksalt crystal symmetry.The depth dependence of defect densities measured for these alloys provided additional advantages: (1) the magnitude of defect segregation to the free surface vs. alloy composition and (2) the bulk defect densities independent of that surface segregation.A combination of depth-dependent and lattice structural techniques over this extended alloy series revealed that native defect densities rather than piezoelectric fields play a role in their near-surface segregation, that Mg in ZnO dramatically reduces native defect densities, that these defect densities depend sensitively on variations in (Mg,Zn)O alloy lattice dimensions, and that the electrostatic energy associated with surfaces and lattice unit cell dimensions can be a key factor in carrier transport and doping of these oxide semiconductors.
We grew five 1 µm thick Mg x Zn 1−x O films with varying Mg concentration on r-plane sapphire.Alloy compositions measured directly from Rutherford backscattering spectrometry (RBS) were x = 0, 0.31, 0.44, 0.52, and 0.56.Using ion channeling experiments (RBS/C), we determined that ∼93% of all Mg incorporated homogeneously in the epilayers occupied Zn sites in the wurtzite structure.XRD showed these films to be single phase, high quality, all with the wurtzite structure and with no cubic inclusions. 22Atomic force microscopy (AFM) showed these surfaces to be smooth on a nanometer scale with no surface asperities.We performed DRCLS measurements using incident electron beam energies E B = 0.1-5 keV from a glancing electron gun in an ultrahigh vacuum (UHV) system with an optical train consisting of a CaF 2 focusing lens, a sapphire viewport, and f -number matcher coupled to an Oriel monochromator and a CCD detector.4][25] Depth dependence of electron-hole excitation was modeled using Monte Carlo simulations. 26For E B = 1, 2, 3, 4, and 5 keV, excitation peaks at depths U 0 = 7, 18, 32, 50, and 72 nm, respectively, with Bohr-Bethe maximum range R B ∼ 3x longer. 138][29] The increase or decrease in work function measured using an atomic force microscope in Kelvin Probe Force Microscopy (KPFM) mode 13 indicates the valence (E V ) or conduction (E C ) band nature, respectively, of such transitions and hence their energy level position within the bandgap.All these measurements were compared with deep level optical spectroscopy (DLOS), current-voltage (I-V), Schottky barrier, and steady state photo-capacitance (SSPC) measurements obtained previously on the same specimens. 30igures 1(a)-1(e) show representative DRCL spectra for all five Mg concentration samples.In each spectrum, the NBE peaks at ≥3.33 eV are the dominant features, increasing in energy with increasing Mg%.Below the NBE are emissions corresponding to defect levels within the band gaps.We subtracted out the second-order replicas of the NBE peak to avoid overlap with defect features.Deep level emission intensities are normalized by NBE intensity to factor out possible variations in  1(a) shows that the 1.77 eV V Zn intensity is approximately 10x lower than the NBE peak intensity.V Zn clustering broadens this energy from 1.7 to 2.0 eV with increasing cluster size. 31,32ith the addition of 31% Mg in Fig. 1(b), this V C intensity drops by nearly an order of magnitude.Similarly, V O intensity decreases to 0.01x of the NBE intensity.Both reach minima for 44% Mg (Fig. 1(c)), then rise gradually for 52% (Fig. 1(d)) and 56% (Fig. 1(e)).Similar defect intensity decreases are evident in MgZnO grown by vapor transport but over a much smaller alloy range. 33,34NBE energies in Fig. 1 increase linearly with Mg content up to x = 0.52, consistent with theory 35 and other reports, [36][37][38] deviating upward for x = 0.56, near the crossover from wurtzite to rocksalt structure. 13igure 1(f) presents characteristic V C and V O intensity profiles with depth showing strong defect segregation toward the surface over tens of nanometers with a minimum at 44% Mg (inset).SPS spectra provided defect level positions in the bandgap for each of the alloys.The contact potential (cpd), i.e., work function difference between the reference AFM probe tip and the (Mg,Zn)O surface, indicates how the Fermi level E F varies with photo-induced population or depopulation of states and thereby band bending within the surface space charge region.For a representative x = 0.31 spectrum in Figure 2, onsets of photostimulated depopulation (n-type positive slope change) from a gap state to E C are evident at 1.85 and 2.5 eV, as with the 3.6-4 eV NBE transition, above which additional free carriers decrease band bending.Population transitions from E V into a gap state (n-type negative slope change) are evident at 2.05 and 2.25 eV.Since their sum nearly equals the bandgap, the 1.85 and 2.05 eV features correspond, respectively, to photo-depopulation and population of the same gap state.Slope changes at 3.6 and 3.86 eV indicate two additional states.Five similar features are evident for all samples, in reasonable agreement with the transition energies of five DLOS trap states reported previously. 30igure 3 shows band gap position of the dominant defect transitions.Here, the ZnO E C is taken as 4.6 eV below the vacuum level, consistent with the electron affinity of the ZnO (10 10) surface, 39 and E C (E V ) increases (decreases) with Mg% following a 2/3-1/3 rule.As Mg content varies, the defect associated with the 2.3 eV ZnO V O level moves nearly parallel with E V , while the 1.77 eV V C level in ZnO tracks with E C .as expected from previous DLOS/t-SPS comparisons. 41,42Total DLOS deep level concentrations for x = 0%, 31%, 44%, 52%, and 56% of 50.6, 6.1, 6.8, 7.8, and 5.4 × 10 16 cm −3 , respectively, 30 also correlate with Figure 4(a).The bulk I(V C )/I(NBE) values for ZnO are consistent with positron annihilation spectroscopy (PAS) calibration values corresponding to ∼0.08 × 10 17 cm −3 in the bulk and 0.25 × 10 17 cm −3 at the surface. 31Previous electrical measurements on these samples also displayed pronounced Schottky barrier decreases and sheet resistance increases above this Mg%, 30 consistent with additional donors.Figure 4(b) shows the XRD-measured variation in (Mg,Zn)O lattice parameter vs. Mg content measured by RBS.Both a-lattice and c-lattice parameters exhibit pronounced minima at x ∼ 0.52. 9A smaller range of alloy composition would not have revealed these XRD and DRCLS minima.
Both electrostatic and thermodynamic factors may contribute to the decrease in V C and V O defect densities with unit cell volume.Electrostatic repulsion may contribute to the free energy associated with defect formation.Thus, V O sites in ZnO result in neighboring Zn atoms with extra charge that would otherwise lead to lattice expansion.Reduction of the unit cell dimension should increase the energy required to form such defects, lowering their density.Analogous effects are reported for native point defects in complex oxides. 43,44The strong decrease in V C density with Mg content may also be thermodynamically driven given the higher bond strength of MgO vs. ZnO, i.e., −∆H 298 (kJ/mol) = 601.6 (MgO) vs. 350.5 (ZnO).Since Mg is energetically more favorable than Zn in filling vacant Zn sites during growth, increasing Mg content would promote decreasing V C density.Density functional theory calculations based on the pressure dependence of defect formation energies are also consistent with the defect density variations in Figure 4(a). 45hese results show that Mg in MBE-grown a-plane ZnO strongly reduces V Zn and V O native point defects, which are mid-gap defects that dominate recombination and follow band edges.Nearly, the same minima of V Zn and V O defect densities in (Mg,Zn)O coincide with minima of their unit cell volumes.This correlation is consistent with the effect of these Zn and O vacancies to increase lattice electrostatic repulsion, thereby increasing formation energies and decreasing their densities.This work reveals a coupling between electronic defect and lattice structural changes and shows that the free energy associated with surfaces, interfaces, and lattice unit cell dimension can be a major factor in carrier transport and doping of these oxide semiconductors.

FIG. 1 .
FIG. 1. (a)-(e) DRCL spectra of Mg x Zn 1−x O for x = 0, 0.31, 0.44, 0.52, and 0.56 and E B = 0.5-4 keV.NBE peak energies increase with Mg, while NBE-normalized deep level emission intensities decrease to a minimum at x = 0.44.(f) Representative segregation profile and bulk threshold for Mg 0.31 Zn 0.69 O and bulk threshold versus Mg% for all samples.

FIG. 2 .
FIG. 2. SPS cpd vs. incident photon energy for Mg 0.31 Zn 0.69 O showing changes in slope at onsets of photo-population and depopulation.

FIG. 3 .
FIG. 3. SPS-derived energy levels of V C and V O within the (Mg,Zn)O band gap vs. Mg content.These midgap V O (V C ) states appear to vary with valence (conduction) bands.

Figure 4
Figure4shows defect densities vs. composition and comparison with lattice constant variation.In Fig.4(a), both V C and V O defect intensities I(V C ) and I(V O ) are normalized to the NBE intensity I(NBE) vs. Mg alloy content.We used DRCL spectra at 2 keV in order to avoid the nearsurface segregated defects, which could increase defect intensities by more than an order of magnitude.Both I(V C )/I(NBE) and I(V O )/I(NBE) exhibit clear minima at x ∼ 0.44.I(V C )/I(NBE) decreases by >100x, while I(V O )/I(NBE) decreases by >30x.Intensity differences measured from two points on the same surface correspond to <10% for x = 0 and <1% for x > 0. Trap state densities measured by t-SPS correlate with DRCLS intensities and display relatively good agreement40 with Gür et al.DLOS results,30 as expected from previous DLOS/t-SPS comparisons.41,42Total DLOS deep level concentrations for x = 0%, 31%, 44%, 52%, and 56% of 50.6, 6.1, 6.8, 7.8, and 5.4 × 10 16 cm −3 , respectively, 30 also correlate with Figure4(a).The bulk I(V C )/I(NBE) values for ZnO are consistent with positron annihilation spectroscopy (PAS) calibration values corresponding to ∼0.08 × 10 17 cm −3 in the bulk and 0.25 × 10 17 cm −3 at the surface.31Previous electrical measurements on these samples also displayed pronounced Schottky barrier decreases and sheet resistance increases above this Mg%,30 consistent with additional donors.Figure4(b) shows the XRD-measured variation in (Mg,Zn)O lattice parameter vs. Mg content measured by RBS.Both a-lattice and c-lattice parameters exhibit pronounced minima at x ∼ 0.52.9A smaller range of alloy composition would not have revealed these XRD and DRCLS minima.Both electrostatic and thermodynamic factors may contribute to the decrease in V C and V O defect densities with unit cell volume.Electrostatic repulsion may contribute to the free energy associated with defect formation.Thus, V O sites in ZnO result in neighboring Zn atoms with extra charge that would otherwise lead to lattice expansion.Reduction of the unit cell dimension should increase the energy required to form such defects, lowering their density.Analogous effects are reported for native point defects in complex oxides.43,44The strong decrease in V C density with Mg content may also be thermodynamically driven given the higher bond strength of MgO vs. ZnO, i.e., −∆H 298 (kJ/mol) = 601.6 (MgO) vs. 350.5 (ZnO).Since Mg is energetically more favorable than Zn in filling vacant Zn sites during growth, increasing Mg content would promote decreasing V C density.Density functional theory calculations based on the pressure dependence of defect formation energies are also consistent with the defect density variations in Figure4(a).45These results show that Mg in MBE-grown a-plane ZnO strongly reduces V Zn and V O native point defects, which are mid-gap defects that dominate recombination and follow band edges.Nearly, the same minima of V Zn and V O defect densities in (Mg,Zn)O coincide with minima of their unit cell volumes.This correlation is consistent with the effect of these Zn and O vacancies to increase lattice electrostatic repulsion, thereby increasing formation energies and decreasing their densities.This work reveals a coupling between electronic defect and lattice structural changes and shows that the free energy associated with surfaces, interfaces, and lattice unit cell dimension can be a major factor in carrier transport and doping of these oxide semiconductors.

FIG. 4 .
FIG. 4. V C and V O defect emission intensities (a) and lattice parameters (b) vs. Mg content.Point-to-point DRCLS variations signified by error bars are smaller than symbols for x > 0.