In this work, we employed X-ray photoelectron spectroscopy to determine the band offsets and interface Fermi level at the heterojunction formed by stoichiometric silicon nitride deposited on AlxGa1-xN (of varying Al composition “x”) via low pressure chemical vapor deposition. Silicon nitride is found to form a type II staggered band alignment with AlGaN for all Al compositions (0 ≤ x ≤ 1) and present an electron barrier into AlGaN even at higher Al compositions, where Eg(AlGaN) > Eg(Si3N4). Further, no band bending is observed in AlGaN for x ≤ 0.6 and a reduced band bending (by ∼1 eV in comparison to that at free surface) is observed for x > 0.6. The Fermi level in silicon nitride is found to be at 3 eV with respect to its valence band, which is likely due to silicon (≡Si0/−1) dangling bonds. The presence of band bending for x > 0.6 is seen as a likely consequence of Fermi level alignment at Si3N4/AlGaN hetero-interface and not due to interface states. Photoelectron spectroscopy results are corroborated by current-voltage-temperature and capacitance-voltage measurements. A shift in the interface Fermi level (before band bending at equilibrium) from the conduction band in Si3N4/n-GaN to the valence band in Si3N4/p-GaN is observed, which strongly indicates a reduction in mid-gap interface states. Hence, stoichiometric silicon nitride is found to be a feasible passivation and dielectric insulation material for AlGaN at any composition.

The AlGaN material system is of high interest in power electronics due to its high dielectric strength that increases with Al composition. This leads to a high theoretical Baliga's figure of merit (BFOM), suggesting the possibility of high power switches with low on-resistances and high breakdown voltages.1–5 However, current switching performance of AlGaN-based devices is away from the theoretical limit.1,6 This is limited by both the surface and bulk defect states. The bulk performance is limited by the compensation at low doping regimes due to the wide bandgap of this material system.7,8 In addition to the compensation, DX center formation in Al rich AlGaN also causes a high dopant activation energy and, therefore, a lower BFOM.9,10 On the surface, lower Schottky barrier heights in GaN (∼1 eV) due to Fermi level pinning11 are a prime challenge to overcome in vertical power Schottky diodes. A similar Fermi level pinning is seen in AlGaN as well12 which impedes ohmic contact formation on Al rich AlGaN.13,14 Further, mid-gap surface states may impede switching performance and cause current collapse.15–18 Consequently, passivation, i.e., reduction in surface/interface state density, is important to produce larger barrier heights, lower leakage currents, and better switching performance. In general, passivation of wide bandgap semiconductors is an integral part of all device structures where control of surface charges and fields is important.

Band offsets between semiconductors and dielectrics are important performance parameters. Typical surface terminations in power electronics, including field plates, require an insulating dielectric that removes any surface charge shielding the field and reduces field crowding at the contact edges. The dielectrics must provide a sufficient electron barrier to prevent carrier injection into the active regions. Such a condition is normally satisfied for narrow bandgap semiconductors. Passivation studies have been reported on GaN and low Al composition AlGaN with various dielectrics.19 However, passivation on Al rich AlGaN is less studied and with a wide bandgap comparable to several dielectrics, there is a probability that the dielectric-semiconductor heterojunction may provide a leakage step into AlGaN. Hence, dielectric-AlGaN interface studies are crucial at high Al compositions to ensure the presence of appropriate band offsets and electron barriers.

Silicon nitride has emerged as the standard passivation material for GaN and AlxGa1-xN(x < 0.3)16–18,20 while exhibiting a relatively large bandgap of 5.3 eV (when stoichiometric).21 It is an excellent diffusion barrier having an amorphous phase with high temperature stability where crystallites are reported to form only at temperatures >1600 K.22 Further, formation of a thin crystalline phase near the interface of low pressure chemical vapor deposition (LPCVD) silicon nitride and AlGaN has been reported, and this may also play a role in passivation.23 However, the bandgap of silicon nitride is lower than AlGaN for x > 0.75 and the feasibility of silicon nitride passivation in Al rich AlGaN and AlN is uncertain. Hence, understanding the nature of band offsets between Si3N4 and AlGaN (x > 0.75) is crucial for determining the feasibility of silicon nitride passivation in Al rich AlGaN.

In this work, we determine the band offsets and electron/hole barriers presented by silicon nitride to AlGaN as a function of Al composition. The reduction in surface and interface states is studied by surface and interface Fermi level measurements, providing information about Fermi level pinning and interface state energies.

A metalorganic chemical vapor deposition reactor (MOCVD) was employed for the growth of the AlGaN epitaxial films on c-plane sapphire used in this work. The metal-polar samples are unintentionally n-type unless specified otherwise. If intentionally doped, Si and Mg were employed as n- and p-type dopants, respectively. N-polar GaN was unintentionally doped with oxygen. All AlGaN films were approximately 300 to 400 nm. Further details on the AlGaN growth process are provided elsewhere.8,24,25 AlGaN was either used as deposited or cleaned with ammonium hydroxide to reduce surface oxide.26 Silicon nitride was deposited using low-pressure chemical vapor deposition (LPCVD) at 800 °C and 300 mTorr with dichlorosilane and ammonia as precursors for Si and N, respectively. The ratio of dichlorosilane and ammonia (40 sccm/120 sccm) was tuned to obtain stoichiometric silicon nitride.

X-ray photoelectron spectroscopy (XPS) was the primary characterization tool used to determine electronic band alignment and the surface/interface Fermi level. It utilizes a dual anode X-ray source with Al (1486.6 eV) and Mg (1253.6 eV) anodes in a UHV chamber with a base pressure of 10−10 Torr and a concentric hemispherical analyzer. XPS calibration and charge correction is detailed elsewhere.11 XPS binding energy measurement resolution is ∼0.2 eV. Ga, Si, C, and valence band were analyzed with Al anode. Due to overlapping of Ga Auger peaks with photo-electron peaks of N and O while using Al anode, N and O were analyzed with Mg anode. The composition of the AlGaN films was determined using an X-ray diffraction (XRD) technique described by Tweedie et al.27 For the XRD measurements, a Philips X'Pert Materials Research Diffractometer with a Cu Kα X-ray source was used. Woollam variable angle spectroscopic ellipsometer was used to determine the optical properties of silicon nitride. Atomic force microscopy (Asylum Research MFP-3D) was used to study the film morphology. Electrical measurements were performed using a Keithley 4200 semiconductor characterization system.

The stoichiometry of the deposited silicon nitride was determined via spectroscopic ellipsometry and XPS chemical analysis. The refractive index is known to be a function of nitride stoichiometry (s = Si/N).28 Amorphous silicon nitride typically varies from stoichiometric (s = 0.75) to silicon rich (s > 0.75) phases. Refractive index is minimum at s = 0.75 (n ∼ 2.03, λ = 632.8 nm) and increases with the incorporation of the α-Si phase.28 The measured refractive index is ∼2.02 at 632 nm indicating stoichiometric Si3N4 deposition. Further, XPS was used to determine the stoichiometry by measuring Si, N, and O content from Si 2p, N 1s, and O 1s core level intensities, respectively, normalized by appropriate atomic sensitivity factors.29 This analysis reveals the presence of Si-N and Si-O bonds as shown in Figure 1 with bulk stoichiometric nitride ((Si-N)/N ∼ 0.77) and surface oxide ((Si-O)/O ∼0.49). Hydrogen content in silicon nitride is independent of Si/N ratio with Si-H/N-H∼1 for XPS measured stoichiometry.30 However, the hydrogen content in LPCVD silicon nitride grown at a relatively high temperature of 850 °C is expected to be <5% (Ref. 31) and no lower binding energy core level peak corresponding to Si-H bond was observed.32 No indications of crystalline phases were observed. AFM surface morphology reveals a uniform conformal deposition of silicon nitride on Al/GaN via step-flow growth mode as shown in Figure 1.

FIG. 1.

(a) Silicon core level (Si 2p) peaks at 101.9 and 103 eV corresponding to the Si nitride and oxide bonds, respectively. AFM characterization (b) and (c) showing uniform conformal deposition of silicon nitride on GaN.

FIG. 1.

(a) Silicon core level (Si 2p) peaks at 101.9 and 103 eV corresponding to the Si nitride and oxide bonds, respectively. AFM characterization (b) and (c) showing uniform conformal deposition of silicon nitride on GaN.

Close modal

Band offset determination via XPS requires measurements of core level and valence band maximum binding energies at the surface of thick films of silicon nitride and AlGaN and at the interface formed by deposition of very thin silicon nitride (∼3 to 4 nm) on AlGaN. Further details can be found elsewhere.33 The valence band offset is determined using the following equation:

ΔEV=[(Ecli)L(EclbEvb)L][(Ecli)R(EclbEvb)R],
(1)

where Ecl and Ev represent core level and valence band maximum energies, respectively. The superscripts i and b represent interface and thick film surface properties, respectively. Si 2p, Si 2s, Al 2p, and Ga 3d core levels were utilized in determining band offsets. XPS valence band analysis (Figure 2) at the surface of thick silicon nitride (155 nm) shows the valence band maximum at 98.9 eV from Si 2p core level. The procedure for determining instrument broadening using standard Au 4f7/2 signal is described elsewhere.11 The valence band maximum in AlGaN is found to be at 70 + 0.8(1-x) (eV) and 17 − 0.8x from Al 2p and Ga 3d core levels, respectively, where x is Al mole fraction. Characterization of valence band maximum and core levels in AlGaN is described and published elsewhere.12 Finally, valence band offsets were determined after determining the core level binding energies at the Si3N4 (3–4 nm)/AlGaN interface as a function of Al composition. The conduction band offset was determined using the known material bandgaps. Figure 3 shows the measured electronic band alignment of silicon nitride on AlGaN as a function of Al composition to be a Type II staggered band alignment with silicon nitride providing an electron barrier via the conduction band offset. The electron barrier and conduction band offset show no significant dependence on Al composition for x < 0.6 and are around 2 eV. The measured conduction band and valence band offsets on GaN are comparable to previous reports with plasma nitridation of silicon on GaN.21 The electron barrier and conduction band offset decrease as Al composition is increased further (x > 0.6) and reduce to 0.7 eV for pure AlN. Si doped GaN (n = 5 × 1016 cm−3) and Al0.6Ga0.4 N (n = 5 × 1017 cm−3) were also analyzed. The band offsets did not change with Si doping, as expected. Interestingly, silicon nitride still provides an electron barrier for AlN (Eg = 6.03 eV) and might be a feasible insulating dielectric even with its lower bandgap (Eg = 5.3 eV). The band offsets (Figure 3) may be given by the following empirical expressions: ΔEC=2+1.5x2.8x2 and ΔEV=0.1+3.4x2.1x2  eV.

FIG. 2.

XPS valence band data from the surface of thick silicon nitride film showing the valence band maximum at 3 eV with respect to the surface Fermi level.

FIG. 2.

XPS valence band data from the surface of thick silicon nitride film showing the valence band maximum at 3 eV with respect to the surface Fermi level.

Close modal
FIG. 3.

The (a) electronic band alignment and (b) band offsets at silicon nitride/AlGaN interface with different Al compositions.

FIG. 3.

The (a) electronic band alignment and (b) band offsets at silicon nitride/AlGaN interface with different Al compositions.

Close modal

In XPS, all binding energies are measured with respect to the Fermi level, i.e., zero binding energy corresponds to EF. Hence, the binding energies of the valence band maxima and core levels also provide the position of the surface/interface Fermi energy within the bandgap. Figure 3 shows the Fermi level (dotted line) at the heterojunction for various AlGaN compositions. Note that the Fermi level shown is only for the Si3N4/AlxGa1-xN interface. The energy difference between the interface Fermi level and the AlGaN conduction band is the electron barrier height into AlGaN at the interface. The Fermi level in the bulk of unintentionally n-type and Si doped samples is assumed to be close to the conduction band. Hence, any band bending (the energy difference between the electron barrier height and bulk Fermi level) is assumed to be upward at the surface when surface Fermi level is pinned lower in the band. Figure 4 shows the measured electron barrier height at the Si3N4/AlGaN interface. Interestingly, a negligible barrier height is observed in AlGaN for x < 0.6. Therefore, the band bending at the surface is also negligible. As with the band offsets, Si doped GaN (n = 5 × 1016 cm−3, EFbulk = EC–0.1 eV) and Al0.6Ga0.4 N (n = 5 × 1017 cm−3, EFbulk = EC–0.05 eV) were also analyzed and the interface Fermi level (and hence the barrier) did not change with the introduction of Si doping. The observed absence of Fermi level pinning and interface state reduction has been observed in HEMTs with AlGaN having Al content <30% (Ref. 15) and our results are consistent with reports on low Al composition AlGaN. Since a large free surface electron barrier height and hence band bending (assumed to be slightly lower than the barrier) on AlGaN of x2+0.8x+0.7  eV with 0.7 eV for GaN, increasing to ∼1.5 eV for Al0.6Ga0.4N12 has been reported (Figure 4), the absence of band bending in our results indicates the feasibility of silicon nitride as passivation dielectric for AlGaN with x < 60%. Further, for x > 0.6, the band bending is reduced by ∼1 eV relative to the free surface due to the interface with silicon nitride.

FIG. 4.

The band bending at the interface of silicon nitride and AlGaN in comparison with that at the free surface as a function of Al molar fraction.

FIG. 4.

The band bending at the interface of silicon nitride and AlGaN in comparison with that at the free surface as a function of Al molar fraction.

Close modal

The observed lack of (x < 0.6) or reduced band bending (x > 0.6) in AlGaN may be understood by analyzing the Fermi level in silicon nitride, which is found to be at 3 eV with respect to the valence band and does not change with AlGaN aluminum composition. We hypothesize that the Fermi level in CVD silicon nitride is determined by the silicon dangling bonds (≡Si0/−1), which produce gap states between ∼3 eV (occupied donor states) and 3.5 eV (unoccupied acceptor states).34,35 A configuration of nitrogen vacancy with hydrogen atom (≡Si-H and ≡Si-Si≡) when occupied may also result in the observed Fermi level at 3 eV. However, unoccupied configuration results in vastly different Fermi level and is less likely as will be discussed later on in this paper. Further, N-H bonds are not expected to be electrically active34,35 and the dominating defect of occupied nitrogen dangling bonds is expected to exhibit a Fermi level at much lower energies.34 It has to be noted that an AlGaN surface subjected to high temperatures under ammonia during silicon nitride deposition may lead to removal of any surface oxygen23 and associated defect states. The significant difference between the Fermi level at the interface and the free surface (with ∼0.5 monolayer of oxygen12) also strongly implies reduced contribution of oxygen associated defect states.

Hence, if the Fermi level of AlGaN is below the Fermi level of Si3N4 before any charge transfer at the heterojunction, electron transfer from silicon nitride to (n-type) AlGaN will produce a (negative) charge accumulation layer that aligns the Fermi level in AlGaN at 3 eV. The conduction band in AlGaN is expected to be close to the Fermi level and is also at approximately 3 eV. The band bending due to the accumulation layer is very sharp and is determined by the Thomas-Fermi screening length ∼aB/2, where aB is the effective Bohr radius (∼1(2) nm in AlN(GaN)) and is independent of charge density.36 This effectively changes the band offsets and explains the nearly constant measured conduction band offsets for AlGaN for x < 0.6 since the conduction band is always found to be aligned to the Fermi level in silicon nitride. For x > 0.6, the conduction band and Fermi level in AlGaN before charge transfer is higher than 3 eV. Thus, the formation of a depletion region in AlGaN to align the Fermi level produces a downward band bending. Band bending in the depletion region extends over hundreds of nm and has no effect on band offsets. This indicates that the observed downward band bending was only a consequence of Fermi level alignment between the silicon nitride and AlGaN. From these results, we could conclude that the Fermi level pinning normally observed on the AlGaN free surface was absent when AlGaN was passivated with stoichiometric silicon nitride. This strongly indicates a reduction in midgap surface/interface states and excellent passivation over the entire composition range of AlGaN. The deep mid-gap states associated with silicon dangling bonds in silicon nitride typically have large time constants (frozen out) and do not respond to the capacitance measurements over typical bias time scales. However, the band alignment at equilibrium is determined by such bulk states. In addition to passivation by silicon nitride, an AlGaN surface subjected to the high temperatures under ammonia during silicon nitride deposition may have led to the removal of any surface oxygen23 and this may also have played a role in the reduction of surface states.

Further analyses via current-voltage-temperature (I-V-T) (Figure 5) and capacitance-voltage (C-V) (Figure 6) characterization of the silicon nitride/GaN heterojunction corroborated the results from the photoelectron spectroscopy. Ni or Ti (diameter of 100 μm for I-V and 500 μm for C-V) were used as electrodes on thin silicon nitride (4 nm) deposited on GaN. The weak temperature dependence of the current in forward bias indicates a (defect assisted) tunneling involving the carriers in the accumulation layer across the silicon nitride. The observed Arrhenius behavior in the reverse bias indicates the expected depletion of the accumulation layer and current limited by Schottky barrier between the electrode (Ni) and GaN and hence supports our hypothesis. Figure 6 shows the C-V measurements on a silicon nitride/(n(5 × 1016 cm−3)/p(2 × 1017 cm−3)-doped) GaN structure. C-V measurements on a Ni-(n-doped GaN) Schottky diode are shown for comparison. The bias-axis intercept showed nearly zero barrier, 1 eV barrier, and barrier comparable to bandgap of GaN for silicon nitride/n-GaN, Ni/n-GaN, and silicon nitride/p-GaN heterostructures, respectively. The shift in the barrier across the bandgap from n-type to p-type GaN indicates highly reduced interface states across the bandgap in GaN and consequently excellent passivation. A similar C-V analysis (Figure 7) of a heterojunction (diameter of 150 μm) with Si3N4 (2 nm) and Si doped Al0.6Ga0.4N (ND-NA = 2 × 1018 cm−3) yielded a low barrier height and band bending of ∼0.3 V in agreement with XPS results (Figure 4). A dielectric capacitance is also seen and it has to be noted that if silicon nitride is the only contributor, the measured relative dielectric constant is low at ∼2 compared to the reported static dielectric constant of 7.4.37 Further, the dielectric is lossy, i.e., for a forward bias >0.5 V, the conductance sharply increases due to the defect assisted tunneling through silicon nitride described earlier with the I-V-T characterization in GaN. Further, C-V measured barrier heights have been found to be unrealistically larger than those measured with XPS and I-V in Schottky diodes on Al rich AlGaN.38,39 I-V curves are also reported to provide large ideality factors.1,39 Hence, the presence of other sources to the dielectric capacitance and barrier inhomogeneity is possible.

FIG. 5.

I-V characteristics of silicon nitride/n-GaN heterostructure.

FIG. 5.

I-V characteristics of silicon nitride/n-GaN heterostructure.

Close modal
FIG. 6.

C-V characterization of (a) silicon nitride/n-GaN and (b) silicon nitride/p-GaN heterostructures.

FIG. 6.

C-V characterization of (a) silicon nitride/n-GaN and (b) silicon nitride/p-GaN heterostructures.

Close modal
FIG. 7.

C-V characterization of silicon nitride (2 nm)/n-Al0.6Ga0.4N heterostructure.

FIG. 7.

C-V characterization of silicon nitride (2 nm)/n-Al0.6Ga0.4N heterostructure.

Close modal

Oppositely charged surface states have been reported to neutralize spontaneous polarization to account for the polarization catastrophe.12,40 After passivation, it is expected that the oppositely charged silicon dangling bond states compensate the spontaneous polarization charge. In order to verify this hypothesis, a heterojunction formed by depositing silicon nitride on N-polar GaN was analyzed where the compensating charge is negative. Similar interface Fermi level and band offsets were observed. Since occupied and unoccupied bands of ≡Si are at ∼3 to 3.5 eV, no significant shift in band offsets or Fermi level is expected if ≡Si were positively or negatively charged as shown in Figure 8. In contrast, configuration of nitrogen vacancy with hydrogen atom (≡Si-H and ≡Si-Si≡) discussed earlier results in vastly different Fermi levels between occupied and unoccupied states and hence is less likely to be responsible. Further, if the compensating charge was in interface states, the oppositely directed interface dipoles would produce different band offsets for nitrogen and metal polar surfaces.11,41,42 Charging of any other defect band in silicon nitride should also have produced a shift in the interface Fermi level. Further, a ≡Si defect density of ∼1020 cm−3 in a 5 nm (5 × 1013 cm−2) thick film would be sufficient to compensate polarization charge (2 × 1013 cm−2(GaN) to 5 × 1013 cm−2(AlN)), i.e., the balancing charge to the bonded polarization charge is expected to be distributed adjacent to the interface in the silicon nitride. The resulting distribution of charge through the silicon nitride is not known, but the direction of interface dipole formed should be similar to the interface states dipole on a free surface reported elsewhere.11 Even slight deviations (<1%) from stoichiometry may have produced the required ≡Si densities. Hence, ≡Si is the likely candidate to compensate the polarization charge.

FIG. 8.

(a) Positive charging of ≡Si0 dangling bonds and (b) negative charging of ≡Si−1 dangling bonds to compensate spontaneous polarization charge (Psp) on GaN.

FIG. 8.

(a) Positive charging of ≡Si0 dangling bonds and (b) negative charging of ≡Si−1 dangling bonds to compensate spontaneous polarization charge (Psp) on GaN.

Close modal

In conclusion, we have determined the band offsets and passivation feasibility of stoichiometric silicon nitride deposited on AlGaN as a function of Al composition. Silicon nitride forms a type II staggered band alignment and presents an electron barrier even at higher Al compositions, where the bandgap of AlGaN is larger than that of silicon nitride. Further, removal of band bending in AlGaN (x < 0.6) and reduction in band bending for x > 0.6 strongly indicate no or greatly reduced surface/interface states. The presence of band bending for x > 0.6 is seen due to Fermi level alignment at the heterojunction and not due to Fermi level pinning by interface states. Photoelectron spectroscopy results were corroborated by C-V measurements. The shift in the barrier across the bandgap for silicon nitride heterojunction with Si doped and Mg doped GaN indicated no midgap interface states and passivation. Hence, stoichiometric silicon nitride presents a feasible passivation material for AlGaN for all compositions.

Partial financial support from NSF (DMR-1108071, DMR-1312582, ECCS-1508854, REU EEC 1156762, and DMR-1508191) and ARO (W911NF-15-2-0068 and W911NF-14-C-0008). We also thank Henry F. Taylor at the NCSU nanofabrication facility (NNF) for his expertise and help with LPCVD silicon nitride deposition.

1.
T.
Kinoshita
,
T.
Nagashima
,
T.
Obata
,
S.
Takashima
,
R.
Yamamoto
,
R.
Togashi
,
Y.
Kumagai
,
R.
Schlesser
,
R.
Collazo
,
A.
Koukitu
, and
Z.
Sitar
,
Appl. Phys. Express
8
,
061003
(
2015
).
2.
I. C.
Kizilyalli
,
A. P.
Edwards
,
H.
Nie
,
D.
Disney
, and
D.
Bour
,
IEEE Trans. Electron Devices
60
,
3067
(
2013
).
3.
I. H.
Oğuzman
,
E.
Bellotti
,
K. F.
Brennan
,
J.
Kolník
,
R.
Wang
, and
P. P.
Ruden
,
J. Appl. Phys.
81
,
7827
(
1997
).
4.
T. L.
Chu
and
R. W.
Kelm
,
J. Electrochem. Soc.
122
,
995
(
1975
).
5.
A.
Raman
,
S.
Dasgupta
,
S.
Rajan
,
J. S.
Speck
, and
U. K.
Mishra
,
Jpn. J. Appl. Phys., Part 1
47
,
3359
(
2008
).
6.
M.
Zhu
,
B.
Song
,
M.
Qi
,
Z.
Hu
,
K.
Nomoto
,
X.
Yan
,
Y.
Cao
,
W.
Johnson
,
E.
Kohn
,
D.
Jena
, and
H. G.
Xing
,
IEEE Electron Device Lett.
36
,
375
(
2015
).
7.
Z.
Bryan
,
I.
Bryan
,
B. E.
Gaddy
,
P.
Reddy
,
L.
Hussey
,
M.
Bobea
,
W.
Guo
,
M.
Hoffmann
,
R.
Kirste
,
J.
Tweedie
,
M.
Gerhold
,
D. L.
Irving
,
Z.
Sitar
, and
R.
Collazo
,
Appl. Phys. Lett.
105
,
222101
(
2014
).
8.
S.
Mita
,
R.
Collazo
,
A.
Rice
,
R. F.
Dalmau
, and
Z.
Sitar
,
J. Appl. Phys.
104
,
013521
(
2008
).
9.
L.
Gordon
,
J. L.
Lyons
,
A.
Janotti
, and
C. G.
Van de Walle
,
Phys. Rev. B
89
,
085204
(
2014
).
10.
R.
Collazo
,
S.
Mita
,
J.
Xie
,
A.
Rice
,
J.
Tweedie
,
R.
Dalmau
, and
Z.
Sitar
,
Phys. Status Solidi C
8
,
2031
(
2011
).
11.
P.
Reddy
,
I.
Bryan
,
Z.
Bryan
,
W.
Guo
,
L.
Hussey
,
R.
Collazo
, and
Z.
Sitar
,
J. Appl. Phys.
116
,
123701
(
2014
).
12.
P.
Reddy
,
I.
Bryan
,
Z.
Bryan
,
J.
Tweedie
,
S.
Washiyama
,
R.
Kirste
,
S.
Mita
,
R.
Collazo
, and
Z.
Sitar
,
Appl. Phys. Lett.
107
,
091603
(
2015
).
13.
R.
France
,
T.
Xu
,
P.
Chen
,
R.
Chandrasekaran
, and
T. D.
Moustakas
,
Appl. Phys. Lett.
90
,
062115
(
2007
).
14.
B. B.
Haidet
,
I.
Bryan
,
P.
Reddy
,
Z.
Bryan
,
R.
Collazo
, and
Z.
Sitar
,
J. Appl. Phys.
117
,
245702
(
2015
).
15.
T.
Mizutani
,
Y.
Ohno
,
M.
Akita
,
S.
Kishimoto
, and
K.
Maezawa
,
IEEE Trans. Electron Devices
50
,
2015
(
2003
).
16.
B.
Bakeroot
,
S.
You
,
T.-L.
Wu
,
J.
Hu
,
M. V.
Hove
,
B. D.
Jaeger
,
K.
Geens
,
S.
Stoffels
, and
S.
Decoutere
,
J. Appl. Phys.
116
,
134506
(
2014
).
17.
B. M.
Green
,
K. K.
Chu
,
E. M.
Chumbes
,
J. A.
Smart
,
J. R.
Shealy
, and
L. F.
Eastman
,
IEEE Electron Device Lett.
21
,
268
(
2000
).
18.
A.
Alexewicz
,
M.
Alomari
,
D.
Maier
,
H.
Behmenburg
,
C.
Giesen
,
M.
Heuken
,
D.
Pogany
,
E.
Kohn
, and
G.
Strasser
,
Solid-State Electron.
89
,
207
(
2013
).
19.
B. S.
Eller
,
J.
Yang
, and
R. J.
Nemanich
,
J. Vac. Sci. Technol. A
31
,
050807
(
2013
).
20.
J. R.
Shealy
,
T. R.
Prunty
,
E. M.
Chumbes
, and
B. K.
Ridley
,
J. Cryst. Growth
250
,
7
(
2003
).
21.
T. E.
Cook
, Jr.
,
C. C.
Fulton
,
W. J.
Mecouch
,
R. F.
Davis
,
G.
Lucovsky
, and
R. J.
Nemanich
,
J. Appl. Phys.
94
,
3949
(
2003
).
22.
R.
Riedel
and
M.
Seher
,
J. Eur. Ceram. Soc.
7
,
21
(
1991
).
23.
X.
Wang
,
S.
Huang
,
Y.
Zheng
,
K.
Wei
,
X.
Chen
,
G.
Liu
,
T.
Yuan
,
W.
Luo
,
L.
Pang
,
H.
Jiang
,
J.
Li
,
C.
Zhao
,
H.
Zhang
, and
X.
Liu
,
IEEE Electron Device Lett.
36
,
666
(
2015
).
24.
R.
Dalmau
,
B.
Moody
,
R.
Schlesser
,
S.
Mita
,
J.
Xie
,
M.
Feneberg
,
B.
Neuschl
,
K.
Thonke
,
R.
Collazo
,
A.
Rice
,
J.
Tweedie
, and
Z.
Sitar
,
J. Electrochem. Soc.
158
,
H530
(
2011
).
25.
A.
Rice
,
R.
Collazo
,
J.
Tweedie
,
R.
Dalmau
,
S.
Mita
,
J.
Xie
, and
Z.
Sitar
,
J. Appl. Phys.
108
,
043510
(
2010
).
26.
S. W.
King
,
J. P.
Barnak
,
M. D.
Bremser
,
K. M.
Tracy
,
C.
Ronning
,
R. F.
Davis
, and
R. J.
Nemanich
,
J. Appl. Phys.
84
,
5248
(
1998
).
27.
J.
Tweedie
,
R.
Collazo
,
A.
Rice
,
J.
Xie
,
S.
Mita
,
R.
Dalmau
, and
Z.
Sitar
,
J. Appl. Phys.
108
,
043526
(
2010
).
28.
J. G. E.
Gardeniers
,
H. A. C.
Tilmans
, and
C. C. G.
Visser
,
J. Vac. Sci. Technol. A
14
,
2879
(
1996
).
29.
J. F.
Moulder
,
W. F.
Stickle
,
P. E.
Sobol
, and
K. D.
Bomben
,
Handbook of X-Ray Photoelectron Spectroscopy: A Reference Book of Standard Spectra for Identification and Interpretation of XPS Data
(
Perkin-Elmer Corporation
,
1992
).
30.
W. A. P.
Claassen
,
W. G. J. N.
Valkenburg
,
F. H. P. M.
Habraken
, and
Y.
Tamminga
,
J. Electrochem. Soc.
130
,
2419
(
1983
).
31.
J.
Yota
,
J.
Hander
, and
A. A.
Saleh
,
J. Vac. Sci. Technol. A
18
,
372
(
2000
).
32.
G. B.
Smith
,
D. R.
McKenzie
, and
P. J.
Martin
,
Phys. Status Solidi B
152
,
475
(
1989
).
33.
E. A.
Paisley
,
M.
Brumbach
,
A. A.
Allerman
,
S.
Atcitty
,
A. G.
Baca
,
A. M.
Armstrong
,
R. J.
Kaplar
, and
J. F.
Ihlefeld
,
Appl. Phys. Lett.
107
,
102101
(
2015
).
34.
J.
Robertson
and
M. J.
Powell
,
Appl. Phys. Lett.
44
,
415
(
1984
).
35.
E.
Vianello
,
L.
Perniola
,
P.
Blaise
,
G.
Molas
,
J. P.
Colonna
,
F.
Driussi
,
P.
Palestri
,
D.
Esseni
,
L.
Selmi
,
N.
Rochat
,
C.
Licitra
,
D.
Lafond
,
R.
Kies
,
G.
Reimbold
,
B.
De Salvo
, and
F.
Boulanger
,
Tech. Dig. - IEEE Int. Electron Devices Meet.
2009
,
1
4
.
36.
C.
Wood
and
D.
Jena
,
Polarization Effects in Semiconductors: From Ab Initio Theory to Device Applications
(
Springer Science & Business Media
,
2007
).
37.
K.
Kobayashi
,
A.
Suzuki
, and
K.
Ishikawa
,
Thin Solid Films
550
,
545
(
2014
).
38.
P.
Reddy
,
I.
Bryan
,
Z.
Bryan
,
J.
Tweedie
,
R.
Kirste
,
R.
Collazo
, and
Z.
Sitar
,
J. Appl. Phys.
116
,
194503
(
2014
).
39.
J.
Tweedie
,
R.
Collazo
,
A.
Rice
,
S.
Mita
,
J.
Xie
,
R.-C.
Akouala
, and
Z.
Sitar
,
Phys. Status Solidi C
9
,
584
(
2012
).
40.
A.
Rizzi
and
H.
Lüth
,
Appl. Phys. Lett.
80
,
530
(
2002
).
41.
S.-C.
Lin
,
C.-T.
Kuo
,
X.
Liu
,
L.-Y.
Liang
,
C.-H.
Cheng
,
C.-H.
Lin
,
S.-J.
Tang
,
L.-Y.
Chang
,
C.-H.
Chen
, and
S.
Gwo
,
Appl. Phys. Express
5
,
031003
(
2012
).
42.
W.-C.
Yang
,
B. J.
Rodriguez
,
M.
Park
,
R. J.
Nemanich
,
O.
Ambacher
, and
V.
Cimalla
,
J. Appl. Phys.
94
,
5720
(
2003
).