Crystalline silicon and germanium are both tetrahedrally coordinated semiconduc-tors with indirect band gaps of 1.12 and 0.66 eV, respectively (Table 5-1).
In microelectronics where silicon devices are state-of-the-art, the application of germanium in complementary metal-oxide semiconductors (CMOS) attracts consider-able attention due to its promise of higher channel mobilities as well as low-voltage operation enabled by the narrower band gap. Most studies performed in this field are dealing with germanium in combination with high permittivity (high-κ) materials such as HfO2 in order to replace SiO2 as the gate dielectric for further scaling of the devices.
However, a direct deposition e.g. by atomic layer deposition (ALD) of the dielectric onto the germanium surface results in a very high density of interface states and hence Fermi level pinning which leads to poor CV-characteristics and high leakage currents [119]. Apparently there exists the need for an ultra-thin (in the order of Ǻ) passivation of germanium surfaces. It has become clear that germanium oxide is not a conceivable candidate for this purpose but rather constitutes a major drawback for a proper surface passivation. Contrary to thermally grown SiOx which drastically reduces the surface
56 Passivating a-Si1-xCx films on crystalline silicon and germanium substrates
states of c-Si, GexOy seriously enhances defect recombination at c-Ge surfaces due to its thermal instability and hygroscopic character. The complete removal of the GexOy
prior to the growth or deposition of a passivating medium is therefore a critical issue.
An improvement of the electrical properties of the devices is accomplished e.g. by the presence of a GeON interlayer which is either formed by thermal nitridation of the Ge surface or by deposition of a GeON layer by atomic beam deposition [119]. However, the introduction of nitrogen near the Ge surface in this approach results in a large density of fixed positive charges that again restrain the electrical performance of the device. As the direct deposition of PECVD SiOx onto the germanium surface results in subcutaneous oxidation at the interface, Wang et al. proposed the introduction of a thin silicon layer which protects the germanium surface from oxidation [120]. This silicon passivation can be realized by different means e.g. by annealing of the samples in SiH4+N2 ambient at 400°C [121] or by growth of a silicon epitaxial layer using SiH4 at temperatures of around 500°C [119]. A subsequent partly oxidation of the Si interlayer can then be performed by immersing the samples, for example, in a O3-based solution.
The preconditioning of the germanium surface in these cases is typically done by etching in hydrofluoric acid and rinsing in de-ionized water (which is an etchant to GexOy). A H2 bake at typical temperatures of 600-650°C is also used to obtain clean Ge surfaces. An interesting issue concerning the epitaxially grown silicon is the ob-served film thickness dependence of the CV-characteristics of the devices [119]. TEM studies reveal that for thicker Si layers the epitaxial structure cannot be preserved on germanium surfaces due to the lattice mismatch between c-Si and c-Ge (Table 5-1) and that the stress in the silicon film relaxes through the formation of a regular array of misfit dislocations.
In photovoltaics (PV) the different absorption characteristics of silicon and germa-nium (Fig. 5-1) account for the different fields in which the semiconductors are ap-plied: c-Si solar cells constitute the major part of today´s terrestrial PV systems, whereas c-Ge is investigated for thermophotovoltaic (TPV) applications [122-124] and for highest efficiency triple-junction solar cells based on GaInP, GaInAs with Ge as the bottom cell. In the latter structure, the emitter of the germanium sub-cell is typi-cally passivated by GaAs or GaInP, whereas there is no need for rear side passivation (of the base) for lattice-matched triple-junction cells: The current density generated mainly in the emitter region of the Ge cell by direct absorption of wavelength
<1600 nm already exceeds the current densities produced in each of the other two sub-cells. However, in lattice mismatched (metamorphic) triple cell structures with higher theoretical efficiency potential, the germanium bottom cell may become current limit-ing and therefore an improvement of current generation in the indirect absorption
Passivating a-Si1-xCx films on crystalline silicon and germanium substrates 57
range is necessary. The latter is sought to be accomplished by a lower doped substrate (higher diffusion length of minority carriers) in combination with an effective rear surface passivation [125].
Regarding the passivation of germanium stand-alone cells (TPV), Posthuma et al.
[122, 126] report on a strong reduction of the surface recombination rate of germanium by low temperature deposition (220°C) of hydrogenated amorphous silicon by means of PECVD. Optimized deposition conditions and a preliminary surface conditioning consisting of a H2O dip in combination with an in-situ hydrogen plasma led to effec-tive carrier lifetimes exceeding 500 µs on 2-3×1015 cm-3p-Ge substrate. However the passivation mechanism is not addressed explicitly.
Si and Ge dangling-bonds have been identified to be responsible for the principal defect states in a-Si and a-Ge [127] and at the surfaces of their respective crystalline counterparts [128]. Regarding the mechanism of passivation of germanium surfaces, a direct saturation of Ge dangling-bonds by Si atoms in the case of the thin epitaxially grown silicon film for microelectronic applications is straightforward. This is sup-ported by the fact that the Si/SiOx system on top of the c-Ge surface results only in a small amount of fixed negative charges [119] which indicates that a passivation by a strong induced band bending at the Ge surface can be discarded. The situation is less clear in the case of passivating hydrogenated amorphous films such as a-Si:H. For the passivation of c-Si surfaces, it is generally agreed that a substantial fraction of the Si dangling-bonds is reduced by the formation of Si-H bonds and that the corresponding bonding and anti-bonding states lie outside the energy gap. However, there exists Fig. 5-1: Absorption coefficient and penetration depth as a function of wavelength for c-Si and
c-Ge. Egap refers to the respective indirect band gaps, EГ1 refers to the respective di-rect band transitions (data modified from [97]).
58 Passivating a-Si1-xCx films on crystalline silicon and germanium substrates
strong experimental evidence that silicon and germanium exhibit rather different char-acteristics in terms of hydrogenation. Studying a-Si:H, a-Ge:H and a-SiGex:H alloy films, Paul et al. stated a limited ability of hydrogen for the passivation of Ge dan-gling-bonds resulting in a strongly enhanced gap state density in a-Ge:H compared to a-Si:H [129]. In a-SiGex:H the authors revealed a preferential attachment of H to Si by a factor of ten and observed hydrogen out-diffusion from a-Ge:H already at tempera-tures as low as 150°C. Jin et al. furthermore found indications for hydrogen-related defects (other than dangling-bonds) in strongly hydrogenated a-Ge [127]. These ex-perimental results are supported by first-principles calculations for the interaction between Ge dangling-bonds and hydrogen [130]. The key findings comprise an energy level of the Ge dangling-bonds below the Ge valence band maximum. This implicates that they are always negatively charged. Furthermore interstitial hydrogen atoms in germanium are found to act exclusively as acceptors. The electrostatic repulsion be-tween H- and the negatively charged defects is hence made responsible for the poor Ge dangling-bond passivation by hydrogen.
Table 5-1: Physical properties of crystalline germanium and silicon at 300°C [14, 94, 131].
properties germanium silicon
atom density (cm-3) 4.42×1022 5.02×1022
dielectic constant, ε 16.0 11.9
effective density of states in conduction band, NC (cm-3) 1.04×1019 2.8×1019 effective density of states in valence band, NV (cm-3) 6.0×1018 2.65×1019
electron affinity, χ (V) 4.0 4.05
energy gap (eV) at 300 K 0.66 (indirect)
0.8 (direct)
1.12 (indirect) 3.4 (direct) intrinsic carrier concentration, ni (cm-3) 2.4×1013 1.00×1010
lattice constant (Å) 5.64613 5.43102
mobility (cm2V-1s-1) electrons μn
holes μp
3900 1900
1450 500
Passivating a-Si1-xCx films on crystalline silicon and germanium substrates 59