The development of remote plasma-assisted treatment of semiconductors has helped to drive microelectronics into the nanoscale regime. Dr. Gerry Lucovsky’s pivotal role in this development has made an impact not only in the realization of ultralow defect densities in SiO2/Si transistor junctions but also in the understanding and control of a wide range of semiconductor interfaces. Here, we show how this technique has expanded our understanding of electronic interfaces beyond the Si interface to many different compound semiconductors.

Professor Gerry Lucovsky made major contributions in many fields of solid-state physics. His later works centered on high-K dielectrics for use as gate dielectrics in Si MOSFETs. These studies addressed a key challenge for transistors—the ability to maintain capacitance C = ɛA/d for dielectric constant ɛ = Kɛ0 while shrinking areal dimensions A and dielectric thickness d, yet minimizing quantum tunneling through the ultrathin dielectric layer. Avoiding this microelectronic “roadblock” required developing higher K dielectrics such as Hf oxide and Zr oxide, a key goal of the microelectronic industry. Such oxides presented their own challenges to remove the dielectric/Si interface states. The reduction of interface state densities by orders of magnitude at the SiO2/Si interface required decades to achieve and to develop new physical and analytic techniques. Understanding and controlling defects at the metal and dielectric interfaces with compound semiconductors required similar developments. In both domains, Dr. Lucovsky was able to couple his solid-state knowledge with a combination of ultrahigh vacuum techniques and physical intuition to address both elemental and compound semiconductor interfaces.

A central tool that Gerry Lucovsky used to make these advances was the development of remote plasma treatment of semiconductors, especially Si, in an otherwise UHV environment. The paper of Yasuda et al. showed that a two-step process of first remote plasma-assisted oxygen to clean the semiconductor surface, then to deposit an oxide, and produce a Si/SiO2 interface with a midgap interface defect density in the 1 × 1010 cm−2eV−1 range as measured by capacitance–voltage (C–V) techniques.1 The use of a remote plasma-assisted oxygen (RPAO) treatment to remove adventitious species before deposited oxides with low interface state densities became a central feature in his subsequent publications. This alternative method surpassed the interface defect density methods from other groups to reduce the interface defect densities, relying on just deposition over the native SiO2 outer layer of Si.2 

Many other advances led by Gerry Lucovsky used this remote plasma approach. Among Lucovsky’s many important papers that included the use of RPAO was a comparison of ultrathin SiO2/Si(100) and SiO2/Si(111) interfaces using a combination of remote plasma oxidation and photoemission spectroscopy techniques to examine atomic-scale bonding differences and resultant defect densities.3 This paper recognized that the SiO2/Si interface at metal-oxide-semiconductor (MOS) devices was not atomically abrupt but consisted of one to two monolayers of silicon suboxide. Using remote plasma oxides, Lucovsky et al. explored how departures from ideal bonding at silicon-dielectric interfaces generate electrically active defects that limit performance and reliability.4 Extending this work to Ge, Lucovsky et al. showed how predeposition plasma nitridation passivated interfaces between crystalline-Ge substrates and Hf-based high-K dielectrics.5 RPAO was used for predeposition of many other oxide and silicate compounds as well.6 

Depth-resolved cathodoluminescence spectroscopy (DRCLS)7–9 of RPAO-treated surfaces provided information on the nature and control of defects at a wide range of semiconductor interfaces, including the high-K dielectric overlayers on Si. The near-nanometer depth resolution of DRCLS10 enabled the measurement of gap state defects in ultrathin, plasma-deposited films and their evolution versus annealing treatments.11–13  Figure 1 illustrates the changes in defects within a 50 Å SiO2 on Si versus annealing treatments commonly used to reduce defects within ultrathin SiO2 films and their interfaces. Emissions at ∼0.6–1.1 eV correspond to the defect states within the 1.11 eV Si bandgap. A broad luminescence extending from 1.5 to 3 eV and centered at ∼1.9 eV is associated with the defects within silicon oxide. Annealing at 900 °C strongly reduces this broad 1.9 eV defect emission but not the higher energy Si-related emissions. Post hydrogenation at 400 °C produces an additional decrease in 1.6–2.9 eV emission relative to bulk Si substrate features.11 

FIG. 1.

DRCLS for the same incident electron-beam energy of 2.0 keV. A Ge detector provided results for the 0.7–1.5 eV range. An S-20 photomultiplier provided results in the 1.4–4 eV photon range. For each detector, the relative intensities of the series of sample treatments directly reflect the luminescence efficiency (Ref. 11). Reprinted with permission from Schäfer et al., Appl. Phys. Lett. 73, 791 (1996). Copyright 1996, AIP Publishing LLC.

FIG. 1.

DRCLS for the same incident electron-beam energy of 2.0 keV. A Ge detector provided results for the 0.7–1.5 eV range. An S-20 photomultiplier provided results in the 1.4–4 eV photon range. For each detector, the relative intensities of the series of sample treatments directly reflect the luminescence efficiency (Ref. 11). Reprinted with permission from Schäfer et al., Appl. Phys. Lett. 73, 791 (1996). Copyright 1996, AIP Publishing LLC.

Close modal

Monte Carlo simulations of the secondary electron cascade14 generated during DRCLS studies provide energy-dependent Bohr–Bethe ranges of maximum excitation depth. Figure 2 illustrates the Monte Carlo simulation of electron cascades generated by incident electron-beam energies of range 0.5–1.0 keV. Depending on the incident beam energy EB, electrons in the low keV range scatter and lose energy successively by x-ray generation, plasmon loss, and ultimately, electron-hole pair creation by impact ionization,15 which then scatter within a few angstroms16 and recombine to excite photon emission. These simulations exhibit characteristic peaks at depths U0 and Bohr–Bethe maximum ranges RB. Digitally subtracting out contributions from the backscattered electrons, the resultant deconvolved excitation profiles appear in Fig. 2 for EB in the 0.5–1 keV range. These simulations in the 0.5–1 keV range show that DRCLS can detect defects with nearly 5–10 nm depth resolution in ultrathin films and their interfaces.

FIG. 2.

Monte Carlo simulations for 0.5–1 keV incident electron-beam energies after removing contributions from the backscattered electrons on 12 nm SiO2/50 nm SiO2/6 nm SiO2. The nanometer scale resolution enables detection of defects within ultrathin films and their interfaces.

FIG. 2.

Monte Carlo simulations for 0.5–1 keV incident electron-beam energies after removing contributions from the backscattered electrons on 12 nm SiO2/50 nm SiO2/6 nm SiO2. The nanometer scale resolution enables detection of defects within ultrathin films and their interfaces.

Close modal

Using this near-nanometer scale DRCLS capability, Lucovsky and collaborators studied the use of nanoscale interlayers to further reduce the defects at the Si/SiO2 interfaces.17 They combined RPAO passivation of the Si surface with plasma-assisted nitridation (RPAN). Figure 3 shows the depth variation of the well-known 2.7 eV E2′ SiO2 defect emission,18 whose relative peak intensity exhibits a maximum at 5–10 nm depth, a 4.3 eV higher lying Si conduction-to-valence band which increases with excitation depth into the Si substrate, and a 2.0 eV emission whose intensity peaks at depths between 10 and 15 nm, indicative of a defect centered at the Si/SiO2 interface and related to oxide bonding. DRCLS has also been used to detect defects at the Si/SiO2 interfaces created by ionizing radiation.19 

FIG. 3.

DRCLS measurement of SiO2 ultrathin films and SiO2/Si interface defects on a near-nanometer scale.

FIG. 3.

DRCLS measurement of SiO2 ultrathin films and SiO2/Si interface defects on a near-nanometer scale.

Close modal

Extension of remote oxidation processing (ROP) to compound semiconductors has proven useful to identify the nature of native point defects and to control their metal Schottky barrier contacts. An essential feature of remote plasma oxidation in studying the semiconductor defects is that it does not introduce additional defects of its own. Conventional RF plasma excitation of chemically active gases produces ions in a plume of ionized gas with enough kinetic energy and momentum to introduce lattice defects at the semiconductor surfaces. An inert carrier gas like helium absorbs momentum and heat from a chemically active gas like oxygen, minimizing the generation of defects in the processed semiconductor surface which the ionized gas contacts.

Figure 4(a) shows DRCL spectra of single crystal ZnO before and after ROP treatment of 20% O2 and 80% He treatment for 1 h.20 The relative CL intensity of the 2.5 eV “green” emission, often attributed to oxygen vacancies,21 decreases with increasing depth and electron-beam energy EB. All spectra are normalized to the near band edge (NBE) peak intensity. (b) The integrated peak intensity ratio I(2.5 eV)/I(3.3 eV) shows a 2× increase within 30 nm of the surface. Following a 1 h exposure, this intensity decreases by 50%, indicating partial removal or passivation of this oxygen vacancy.

FIG. 4.

DRCL spectra of ROP-treated and untreated Eagle–Picher ZnO single crystals. (a) 2.5 eV sub-bandgap emission spectra normalized to the 3.3 eV near band edge (NBE) peak intensity, decreasing with increasing incident beam energies and depth into the bulk crystal. (b) Ratio of the normalized “green” defect intensity for this ZnO before vs after the ROP treatment, exhibiting a strong near-surface defect decrease after the ROP treatment. (Ref. 20). Reprinted with permission from Mosbacker et al., Appl. Phys. Lett. 87, 012102 (2005). Copyright 2005, AIP Publishing LLC.

FIG. 4.

DRCL spectra of ROP-treated and untreated Eagle–Picher ZnO single crystals. (a) 2.5 eV sub-bandgap emission spectra normalized to the 3.3 eV near band edge (NBE) peak intensity, decreasing with increasing incident beam energies and depth into the bulk crystal. (b) Ratio of the normalized “green” defect intensity for this ZnO before vs after the ROP treatment, exhibiting a strong near-surface defect decrease after the ROP treatment. (Ref. 20). Reprinted with permission from Mosbacker et al., Appl. Phys. Lett. 87, 012102 (2005). Copyright 2005, AIP Publishing LLC.

Close modal

The removal of oxygen vacancies from ZnO single crystal surfaces produces major changes in metal—ZnO Schottky barriers.22  Figure 5 shows how ROP treatments of the as-received, polished, and etched ZnO single crystal reduces 2.5 eV oxygen vacancy emission in the outer 30 nm, first with 30 min O2 plasma treatment followed by a second ROP exposure which reduces it further. The as-received 2.5 eV normalized DRCLS depth profile exhibits a dominant 2.5 eV peak intensity relative to the 3.3 eV NBE peak intensity—I(2.5 eV)/I(NBE), which decreases strongly with these successive treatments. Figure 5(b) shows the effect of this sequence of plasma treatments on the Au metal/ZnO diode’s rectifying properties.

FIG. 5.

(a) Relative defect intensity changes vs depth and vs ROP exposure. (b) Corresponding I–V plots before and after this ROP treatment, followed by H2 plasma exposure for Au diodes on these surfaces (Ref. 22). Reprinted with permission from Brillson et al., Appl. Phys. Lett. 90, 102116 (2007). Copyright 2007, AIP Publishing LLC.

FIG. 5.

(a) Relative defect intensity changes vs depth and vs ROP exposure. (b) Corresponding I–V plots before and after this ROP treatment, followed by H2 plasma exposure for Au diodes on these surfaces (Ref. 22). Reprinted with permission from Brillson et al., Appl. Phys. Lett. 90, 102116 (2007). Copyright 2007, AIP Publishing LLC.

Close modal

The as-received polished and etched (AR) I–V plot exhibits ohmic behavior with a straight line through the origin. With a 30 min O2/He ROP, the I–V plot exhibits rectifying behavior and a barrier height ΦB = 0.44 eV. With the additional 30 min ROP, the I–V plot exhibits even stronger rectifying behavior with ΦB = 0.51 eV and an ideality factor n = 1.57. These ROP treatments are consistent with oxygen filling the oxygen vacancy sites in the ZnO lattice. Introducing an H2/He remote plasma treatment (RHP) of this diode has the opposite effect, producing I–V ohmic behavior, which is even steeper and more ohmic than before ROP exposure. This striking reversal is consistent with hydrogen removing oxygen atoms from within the outer ZnO surface and increasing oxygen vacancies VO. Likewise, RHP combined with Hall effect and vibrational spectroscopy show that RHP can increase the electron concentration in ZnO single crystals by more than an order of magnitude. These measurements suggest a strong link between the incorporation of hydrogen, increasing donor-bound exciton photoluminescence and increased n-type conductivity.23 This increased defect density can then contribute to trap-assisted tunneling, reduced rectification, and ohmic contact behavior at high enough trap densities.24,25

ROP treatment of the semiconductors also enables direct correlation of defect densities measured by DRCLS with free carrier densities near the surfaces. Figure 6 shows 2 keV (U0 = 17 nm) CL spectra measured through Au diodes on relatively low defect ZnO (0001) Zn-face and ( 000 1 ¯ ) O-face polished and etched single crystals.26 Gold deposited on the as-received ZnO’s Zn and O faces formed the “Au I” diodes shown for each Zn-face and O-face surface. Subsequently, ROP treatment followed by Au deposition of the Zn-face and O-face surfaces with the Au I diodes produced a second set of diodes termed “Au II.” The resultant Au I and Au II diodes on the Zn-face and O-face surfaces enabled a direct comparison of the electric and optical properties for diodes (i) with and without ROP treatment and (ii) comparison of these properties for Zn-face versus O-face surfaces.

FIG. 6.

Comparison of 10 K CL spectra for (a) Au I and (b) II diodes on Zn and O faces. The corresponding CV carrier profiles appear in the insets (Ref. 26). Reprinted with permission from Dong et al., J. Appl. Phys. 108, 103718 (2010). Copyright 2010, AIP Publishing LLC.

FIG. 6.

Comparison of 10 K CL spectra for (a) Au I and (b) II diodes on Zn and O faces. The corresponding CV carrier profiles appear in the insets (Ref. 26). Reprinted with permission from Dong et al., J. Appl. Phys. 108, 103718 (2010). Copyright 2010, AIP Publishing LLC.

Close modal

Figure 6(a) shows CL features corresponding to the ≤3.3 eV ZnO NBE peak and phonon replicas, together with deconvolved gap state emissions at 2.1 and 2.5 eV. The 2.1 eV peak emission corresponds to Zn vacancy (VZn) defect clusters as identified using DRCLS, surface photovoltage spectroscopy (SPS), and positron annihilation spectroscopy (PAS).27 CL spectra in Fig. 6(a) for the Zn-face (in red, labeled Au I) show a broad midgap defect peak for the Au I diode, centered at 2.5 eV, like the 2.5 eV VO peak in Fig. 4(a) but with higher intensity. The CL spectrum for the Au II diode (in green, labeled Au II) shows both the 2.5 and 2.1 eV peaks, indicating the presence of both VO and VZn defects under the Au II diode.

The insert in Fig. 6(a) shows the corresponding C–V measurements for the Zn-face Au I and Au II diodes. The Au I diode shows a free carrier density >0.6 × 1017 cm−3 extending to within 80 nm depth below the metal/ZnO interface, a depth corresponding to its depletion width. The Au II diode shows a similar bulk free carrier density but decreasing from 240 to 140 nm, corresponding to a depletion width nearly 2× larger, indicating a free carrier density ∼4× lower. This difference between Au I and Au II diodes indicates that the additional VZn defects induced by the Au contact on the ZnO Zn-face reduces the free carrier density for the Au II diode vs the Au I diode on the Zn-face. In contrast, the Fig. 6(b) CL spectra for the ZnO O-face shows no 2.1 eV features corresponding to VZn for either Au I or Au II diodes. Furthermore, the inset in Fig. 6(b) shows a higher free carrier density between 80 and 180 nm depths for the Au II diode with no decrease toward the interface. Much higher free carrier densities prevented C–V measurements for the Au I diode, consistent with the higher 2.5 eV VO donor density, and the absence of VZn acceptors.

The comparison of metal diode defect vs electrical properties enabled by ROP demonstrates that Zn- and O-polar ZnO surfaces and Schottky barrier diodes display strong electronic and optical differences, and these differences correlate with defect emissions, traps, and metal/ZnO interface chemistries.28 Even more generally, these Zn-face vs O-face differences are consistent with the major differences between semiconductor anion and cation surface chemistries.29 

Extended from elemental to II–VI compound to III–VI compound semiconductors, ROP processing enables the direct physical identification of a primary native point defect. The ultrawide bandgap semiconductor Ga2O3 has emerged over the past decade as a leading candidate for high power electronics due to its 4.8 eV band and predicted dielectric breakdown field gradient of 8 megavolts.30,31 Native point defects in Ga2O3 are of major concern since they play a dominant role in free carrier density, mobility, and dielectric breakdown.32 Theoretical studies of defects in Ga2O3 suggest several different types of defects including native defects,33 native defect complexes,34 and self-trapped excitons.35 DRCLS is one of the few measurements to physically identify these various Ga2O3 defects.

The DRCLS measurement of defects in Ga2O3 takes advantage of its near-nanometer probe depth control near surfaces. This capability enables the measurement of relatively small volumes of semiconductors whose surfaces can be chemically processed to produce changes which are characteristic of specific defects. Other characterization techniques that provide physical signatures of defects can give significant quantitative information, but have minimum number of defect thresholds requiring physical volumes which are significantly larger than that needed for DRCLS. A notable exception which complements DRCLS and can provide quantitative information is SPS, whose near-surface detection thresholds based on electrostatics are comparable.36,37 Spatial detection of extended bulk defects, e.g., dislocations, are more challenging for all.

DRCLS of Ga2O3 grown by various methods yields characteristic optical spectra which provide evidence for multiple native point defects.38–41, Figure 7(a) shows such characteristic DRCLS spectra of low-pressure chemical vapor deposition (LPCVD)-grown Ga2O3 at EB increasing from 0.5 to 5 keV, corresponding to a depth range from 10 to 100 nm as shown in Fig. 7(b).38 DRCL spectra of edge-fed growth (EFG) and MBE-grown Ga2O3 exhibit similar features.42 Deconvolution of these spectra appears in Fig. 7(a). The ROP treatment of Ga2O3 produces significant changes in these DRCL spectra but which are limited to only one of these peaks, the dominant 3.5 eV feature in LPCVD Ga2O3 and in two sequential ROP exposures. Before ROP exposure, the as-received 3.5 eV feature has a constant intensity vs depth, consistent with a native point defect throughout the bulk. A 1 h 45 min ROP exposure decreases this 3.5 eV area smoothly by 33% from 80 nm to below 10 nm. A subsequent ROP for a total 3 h 15 min exposure decreases the 3.5 eV peak area by 75% from 100 nm to below 10 nm. In comparison, the 2.5 and 3.0 eV peak areas showed no similar decreases. The 3.5 eV peak feature’s steady decrease toward the surface and its deeper penetration with increasing ROP exposure provide evidence for oxygen diffusion into the Ga2O3 lattice and removal of VO lattice sites.

FIG. 7.

(a) DRCL spectra of LPCVD Ga2O3 after 1.75 h ROP with sub-bandgap features, 4.6–4.8 eV bandgap shoulder, and changes with depth. The dashed line outlines Gaussian peak deconvolutions for the EB = 0.5 keV spectrum. (b) Depth profiles of normalized (to NBE) 3.5 eV peak areas vs RB for as-received LPCVD Ga2O3 and two sequential ROP exposures. Analogous 2.5 eV (c) and 3.0 eV (d) profiles before and after 1.75 h ROP exposure show no significant changes (Ref. 38). Reprinted with permission from Gao et al., Appl. Phys. Lett. 112, 242102 (2018). Copyright 2018, AIP Publishing LLC.

FIG. 7.

(a) DRCL spectra of LPCVD Ga2O3 after 1.75 h ROP with sub-bandgap features, 4.6–4.8 eV bandgap shoulder, and changes with depth. The dashed line outlines Gaussian peak deconvolutions for the EB = 0.5 keV spectrum. (b) Depth profiles of normalized (to NBE) 3.5 eV peak areas vs RB for as-received LPCVD Ga2O3 and two sequential ROP exposures. Analogous 2.5 eV (c) and 3.0 eV (d) profiles before and after 1.75 h ROP exposure show no significant changes (Ref. 38). Reprinted with permission from Gao et al., Appl. Phys. Lett. 112, 242102 (2018). Copyright 2018, AIP Publishing LLC.

Close modal

Density functional theory (DFT) calculations of charge state transitions for VO, VGa, and VGa−H provide further support for the identification of the 3.5 eV peak feature with VO. Together with SPS to determine the position of energy levels within the Ga2O3 band gap,40 the combined DRCLS-SPS energies closely match the predicted energy levels for all three defects on their various lattice sites. In particular, the 3.5 eV SPS photopopulation transition and the 3.5 CL transition agree closely with the 3.52 eV VO+(III) charge state energy predicted theoretically and, while not definitive, this configuration is more thermodynamically stable than its VO+(I) and VO+(II) lattice site counterparts.43 

The RPAO technique developed by Gerry Lucovsky has led to numerous advances in the understanding and control of electronically active point defects at SiO2/Si and metal interfaces with II–VI, III–VI, and other semiconductor oxides. Remote oxygen, nitrogen, and hydrogen plasmas together with DRCLS have provided powerful new methods to understand the nature of local chemical bonds on an atomic scale and to control their electronic properties on a macroscopic scale.

This work was sponsored previously in part by the National Science Foundation (NSF) (Grant No. DMR-18-00130), the Office of Naval Research (ONR), the Semiconductor Research Corporation (SRC), and the NSF Engineering Research Center at North Carolina State University.

The author has no conflicts to disclose.

Leonard J. Brillson: Conceptualization (equal); Data curation (equal); Formal analysis (equal); Funding acquisition (equal); Investigation (equal); Methodology (equal); Project administration (equal); Resources (equal); Writing – original draft (equal); Writing – review & editing (equal).

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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