The firing stability of a surface passivation scheme usually refers to the preserva-tion of the passivapreserva-tion quality after a short high temperature step. The latter is neces-sary for the front contact formation of a solar cell when using metal containing pastes deposited onto the (SiNx) anti-reflection coating. The firing process lasts typically only a few seconds and involves peak temperatures between 750-900°C.
A set of silicon lifetime samples passivated by Si-rich a-Si1-xCx layers was subject to comparable temperature loads and the effective lifetimes of the samples were meas-ured before and after the temperature treatment. The prepared passivation schemes comprised a variation in the carbon content of the a-Si1-xCx film (variation of CH4
flow) as well as the capping of the Si-rich film by C-rich films or PECVD a-SiOx. None of the investigated structures could preserve the initial passivation quality owing to the rapid out-diffusion of hydrogen from the amorphous network provoking there-fore a collapse of the prevailing means of passivation. In this respect the investigated Si-rich a-Si1-xCx is more closely related to a-Si:H than to SiNx. The thermal stability at elevated temperatures of the latter is inherently linked to the existence of (temperature stable) fixed charges. Yet no evidence of carbon related charges in the a-Si1-xCx matrix could be found in this work.
As described in section 5.3, the introduction of a small amount of carbon into the matrix impacts the temperature stability in the low temperature range (300-400°C).
This characteristic allows for an increased process temperature tolerance in high-efficiency low temperature solar cell approaches such as heterostructure concepts and is therefore an attractive material for the substitution of a-Si:H.
Likewise, equally prepared germanium samples suffered from a complete depassi-vation after a typical firing process with a peak temperature of 800°C. This finding points to Ge-Si (Ge-C) bond rupture (and reconstruction of the amorphous network) at elevated temperatures. This behavior is in line with the assumption that no charges are involved in the surface passivation of c-Ge by a-Si1-xCx.
5.7 Chapter summary
The main differences in terms of surface passivation of crystalline silicon and ger-manium substrates consist in the instability of (thermally grown) GexOy and the limited practicability of hydrogen for the passivation of Ge dangling bonds. These two major
92 Passivating a-Si1-xCx films on crystalline silicon and germanium substrates
drawbacks are responsible for the preference of silicon to germanium in microelec-tronic technology today.
The performed TEM studies reveal no difference in the film structure depending on the substrate type. However, the feasible resolution was not sufficiently high as to gain insight into the first monolayers at the interface where the impact of the substrate is expected to be pronounced. The results obtained via TEM are in accordance with FT-IR and ellipsometry measurements, giving moreover no evidence on structural differences of the deposited films related to the substrate type.
With increasing C-content, the Si-rich a-Si1-xCx films become less dense which is attributed to an increasing microvoid density. The latter is inferred from an increasing amount of bonded hydrogen present in the matrix. The c-Si surface passivation quality is found to be directly correlated with the evolution of the Si-H bond density in the film. The onset of Si-H bond rupture at around 300°C therefore coincides with the onset of electrical degradation independently of the carbon content and the doping density in the film. The latter, however, has a strong impact on the absolute level of passivation quality. In the case of c-Si substrate, phosphorous doped a-Si1-xCx allows for an effective suppression of surface recombination, boron doping, on the other hand, drastically deteriorates the passivation quality. This observation is addressed to a de-crease in Si-H bond density (implicating an inde-crease in silicon dangling bonds) due to a Fermi-level dependent Si-H bond rupture in the film as well as to an increased re-combination activity due to boron related defects. The surface passivation of germa-nium by a-Si1-xCx demonstrates a fairly different behavior. First of all, the thermal stability of intrinsic films is significantly increased. The electrical degradation starts at temperatures as high as 450°C and therefore indicates a “decoupling” of hydrogen present in the film and passivation quality. Contrary to c-Si substrate, experiments with doped films for the passivation of germanium indicate good electrical properties of p-type and a complete depassivation for n-type a-Si1-xCx.
Regarding the optimum substrate temperature during film deposition, a value of Topt=270°C was found to provide the best overall performance for Si as well as for Ge substrate. A similar behavior of both substrate types is observed for deposition tem-peratures below and above this temperature. Post-deposition annealing improves the passivation quality for Tann<Topt and results in a degradation for Tann >Topt. This finding is in accordance with the role of hydrogen for the passivation of c-Si surfaces. As a result of the experiments, the carbon in the a-Si matrix is supposed to inhibit the epi-taxial growth observed for a-Si depositions at increased temperatures. In the case of c-Ge, the level of surface passivation is assumed to be indirectly correlated with the
Passivating a-Si1-xCx films on crystalline silicon and germanium substrates 93
hydrogen incorporation at increased deposition temperatures since the latter is sup-posed to have a direct impact on the growth kinetics of the film.
From isothermal lifetime experiments, an activation energy for the degradation of the passivation quality was extracted. For the c-Si/a-Si1-xCx system, this energy amounted to approx. 0.5 eV, for the c-Ge/a-Si1-xCx system, a value of approx. 2 eV was found, clearly pointing to different passivation mechanisms involved. Furthermore the results of this type of experiment allow for linking the temperature stability of the electrical properties of the film to the equilibrium between Si-H bond rupture and its inverse process (Si-H↔Si + H). The latter is triggered by the availability of free atomic or molecular hydrogen in the matrix. This simple model is supported by hydro-gen effusion measurements which moreover reveal the impact of carbon on hydrohydro-gen diffusion. Two contrary effects are encountered for Si-rich a-Si1-xCx films for an in-creasing C fraction: on the one hand, the matrix becomes less dense (increased H-diffusivity), on the other hand the increased C-density seems to delay the hydrogen out-diffusion (shift of main peaks to higher temperatures).
A set of passivated silicon wafers of different thicknesses allowed for the extrac-tion of the injecextrac-tion dependent effective SRV of the a-Si1-xCx films. The obtained data do not depend on assumptions concerning the bulk lifetime of the material. The fitting of the SRV data of p- and n-type silicon substrate with the extended SRH model re-veals a change of sign for the fixed charge density parameter Qf of the film when changing the doping polarity of the substrate. The existence of amphoteric interface (near) states is confirmed by SPV measurements. For the passivated germanium sam-ples, only a small surface band bending is measured which does not seem to be altered through annealing. As the latter improves the passivation quality, a field-effect passi-vation of the germanium surface can be discarded. It is therefore concluded that a direct saturation of Ge dangling-bonds by Si- and/or C-atoms is responsible for the suppression of carrier recombination at the surface. The passivation of c-Si surfaces on the contrary is inherently linked to the saturation of Si dangling-bonds by hydrogen.
Although the incorporation of carbon clearly enhances the thermal stability of the c-Si passivation in the temperature range up to 500°C (as compared to pure a-Si), the stability during firing processes (750-900°C) could not be verified.
6 Silicon solar cells with a-Si
1-xC
xrear side schemes
In this chapter the surface passivation properties of Si-rich a-Si1-xCx films are analyzed at the device level. Solar cells with different amorphous silicon carbide rear side schemes are presented. First of all, p-type silicon solar cells with intrinsic, boron and phosphorous doped a-Si1-xCx rear side passi-vation are presented. The application of different rear contacting techniques helps to elucidate the impact of the film doping on the respective passivation mechanism. The combination of doped a-Si1-xCx passivation layers with an adequate laser process results in the so-called PassDop approach. During the laser process, the doped film acts as a dopant source. The induced local back surface fields are characterized in terms of doping profile and electri-cal performance. Applying the PassDop passivation and contacting scheme to n- as well as to p-type solar cells reveals the very high potential of this in-dustrially feasible approach. Finally, a simplified solar cell process based on the rear side passivation by an in-situ diffused back surface field from doped a-Si1-xCx films is presented. This approach is shown to have an efficiency po-tential of up to 20 %.