An approach for the solar cell front side based on amorphous silicon carbide nec-essarily comprises at least two different layers. A thin Si-rich film (< 10 nm) next to the interface is needed for the electrical surface passivation and a C-rich film of de-fined layer thickness assumes the role of the anti-reflection coating (ARC).
Si-rich a-Si1-xCx films exhibit good passivation properties on planar (shiny etched) phosphorous diffused n+-surfaces, however a strongly deteriorated performance is observed on textured (random pyramids) surfaces. The latter is ascribed to a detrimen-tal impact of the pyramids on the film growth. Stack systems consisting of Si-rich a-Si1-xCx(i)/a-Si1-xCx(n) films yield comparable or even better results than merely intrinsic passivation layers. Boron diffused p+-surfaces could not be satisfactorily
Fig. 7-12: EQE and reflection (R) measurements and calculated IQE for bifacially amor-phous silicon carbide coated solar cells with shiny etched front surface (left) and front surface with random pyramid texture (right).Parasitic absorption loss (blue are) and electrical loss (yellow area) are estimated.
140 Amorphous silicon carbide for the solar cell front side
passivated by a-Si1-xCx deposited at the typical temperature of 270°C. Lower deposi-tion temperatures in combinadeposi-tion with subsequent annealing were found to improve the passivation quality. This behavior is attributed to an out-diffusion of hydrogen from the a-Si1-xCx film from interface near regions at an early stage of the deposition process due to the Fermi-level dependent Si-H bond-rupture in the a-Si1-xCx film.
The typical procedure for the extraction of the emitter saturation current (J0e) from lifetime measurements was found to be unfeasible for the surface passivation by a-Si1-xCx. This finding is attributed to the strongly injection dependent surface recom-bination velocity of a-Si1-xCx at low injection levels independently of the surface dop-ing polarity. Hence the prerequisite for the concept of a constant emitter saturation current is not fulfilled and strictly speaking the latter is not valid in the case of a-Si1-xCx passivation. However, in the case of negligible recombination in the base (which is given for high resistivity Fz-silicon material), the measured implied Voc at one sun illumination can be used to estimate the equivalent J0e at the corresponding injection level. Following this approach, J0e values of 40 fA/cm2 for planar and of 400 fA/cm2 for textured (random pyramids) n+-surfaces were found (featuring 10 nm of a-Si1-xCx(i) passivation). The values for shiny etched p+-emitters amount to approx.
900 fA/cm2.
C-rich a-SiyC1-y (y >0.5) layers were optimized regarding their transparency, yield-ing ARC properties slightly inferior to the ones of typical ARC-SiNx. A further im-provement, however, seems to be at hand by adjusting the reactor setup for lower process pressures and higher microwave (peak) powers using pulsed microwave gen-erators.
A front contact formation by metal pastes requiring a firing step is not compatible with the a-Si1-xCx passivation. Therefore a low temperature approach (SE-PassDop) was presented which basically consists of a laser process for opening of the silicon carbide and of subsequent metallization by Ni/Ag or Ni/Cu plating. During the laser process, additional dopant atoms incorporated in the front passivation scheme are driven into the surface and allow for the formation of a selective emitter. An increased laser induced surface damage was found in the case of textured surfaces.
First p-type solar cells featuring the PassDop rear and the SE-PassDop front side approach for the passivation and contacting of respective surfaces were fabricated with planar (shiny etched) and textured (random pyramids) front side. The very low short-circuit currents of the cells are attributed to a strong parasitic absorption on the front side, mainly due to a non-optimized C-rich a-SiyC1-y layer. Open-circuit voltages of up to 672 mV and 623 mV evidence a high level of surface passivation on planar emitters
Amorphous silicon carbide for the solar cell front side 141
and a strongly deteriorated performance on textured surfaces, respectively. The feasi-bility of the metallization concept including the laser process is impressively demon-strated by fill factors of up to 79 % irrespective of the surface condition. The incorpo-ration of a doped, Si-rich a-Si1-xCx film in the emitter passivation scheme proved to be beneficial in terms of series (contact) resistance as well as in terms of front surface recombination.
Finally, it is believed that the parasitic absorption in the front a-SiCx layer stack can be reduced strongly by application of an optimized C-rich a-SiyC1-y layer and further reduction of the thickness of the passivating/dopant containing films. However, it is not yet clear how to overcome the problems related to the poor electrical passiva-tion of surfaces featuring random pyramids.
8 Summary
The amorphous silicon carbide films investigated throughout this work were pro-duced on the basis of plasma enhanced chemical vapor deposition (PECVD). The two different PECVD reactors used for the deposition of the films were presented. The principle development of the material was conducted at the AK400M batch-type reac-tor from Roth&Rau and the processes were then transferred to an industrial type SINA in-line reactor from the same company. For standard passivating Si-rich a-Si1-xCx
layers deposited in the AK400M reactor working with RF power only, SIMS meas-urements revealed rather low C-atom densities (< 10 at.%) in the matrix. Doping atom densities of 1×1021 cm-3 for phosphorous and of 5×1020 cm-3 for boron were deter-mined for a typical gas flow of 100 sccm (B2H6 or PH3 highly diluted in H2). The a-Si1-xCx films fabricated in the SINA reactor exhibit a similar composition as well as similar electrical properties as their batch-type counterparts.
A parallel analysis of a-Si1-xCx/c-Si and a-Si1-xCx/c-Ge systems was performed.
From literature it is known that the main differences in terms of surface passivation of crystalline silicon and germanium substrates are the instability of (thermally grown) GexOy and the limited practicability of hydrogen for the passivation of Ge dangling bonds. From a compositional viewpoint, the performed transmission electron micros-copy (TEM) studies reveal no difference in the film structure depending on the sub-strate type. With increasing C-content, the Si-rich a-Si1-xCx films become less dense, a fact which is attributed to an increasing microvoid density. The c-Si surface passiva-tion quality is found to be directly correlated with 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 surface passivation of germanium 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 there-fore indicates a “decoupling” of hydrogen present in the film and passivation quality.
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 the Ge substrate. A similar behavior for both substrate types is observed for deposition temperatures 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
144 Summary
epitaxial growth observed for a-Si depositions at increased temperatures. The latter is considered to be responsible for the electrical degradation of the film. In the case of c-Ge, the level of surface passivation is assumed to be indirectly correlated with the 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 amounts 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.
A set of a-Si1-xCx passivated silicon wafers of different thicknesses allowed for the extraction of the “pure” injection dependent effective SRVs. The fitting of the SRV data of p- and n-type silicon substrate with the extended Shockley-Read-Hall (SRH) model reveals 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 inter-face (near) states is confirmed by SPV measurements. Several experiments performed with the a-Si1-xCx/c-Ge system point to a direct saturation of Ge dangling-bonds by Si- and/or C-atoms. The latter effectively suppresses carrier recombination at the germa-nium 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 in the a-Si matrix was shown to clearly enhance the thermal stability of the c-Si passi-vation 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.
The performance of intrinsic and doped a-Si1-xCx as rear side passivation (in im-mediate contact to the solar cell base) was investigated. Intrinsic films preserve their excellent passivation quality at the solar cell level, in particular no indication of inver-sion layer shunting comparable to the phenomenon observed in SiNx is found. The performance of its doped counterparts as passivating media strongly depends on the doping level and on the applied rear contacting scheme. However, the major benefit of doped Si-rich a-Si1-xCx layers is unveiled in combination with (thin) intrinsic films.
This is demonstrated by a newly developed rear passivation and contacting scheme based on intrinsic and doped a-Si1-xCx films in conjunction with an adequate laser process (PassDop). The fulfilled demands of the amorphous silicon carbide system on the rear of the solar cell are threefold: very low SRVs are yielded by the intrinsic film
Summary 145
next to the crystalline silicon interface, an overlying highly doped Si-rich a-Si1-xCx
layer acts as dopant source during the laser process and a C-rich layer enhances the optical confinement at the rear of the cell. On high-efficiency n-type solar cells, this approach evidenced a very high process stability and led to efficiencies of up to 22.4 % (Voc =701 mV, FF =80.1 %). On p-type cells, the equivalent process yielded efficien-cies of up to 21.6 % (Voc =683 mV, FF =80.7 %) being therefore directly comparable to passivation schemes featuring high quality thermal SiO2 and to atomic-layer depos-ited Al2O3 in combination with the LFC process.
The use of a highly doped a-Si1-xCx layer as a doping source during the solar cell fabrication can result in a thermally stable high-low passivation scheme. A simplified process sequence for n-type solar cells based on the in-situ diffusion of a n+-BSF from phosphorous doped a-Si1-xCx was introduced. The diffusion of the rear high-low junc-tion takes place during the high-temperature process for the formajunc-tion of the boron emitter. The rear contacting is done by a laser process equal to the one presented in the PassDop approach. The introduction of an additional C-rich film on the rear makes this approach optically superior to diffused full area contacted BSFs. In combination with a well passivated front side, this concept was found to have an efficiency potential of up to 20 %.
Layer structures based on amorphous silicon carbide were also applied to the solar cell front side. The former necessarily comprise at least a two layer system. A thin Si-rich film (≈ 10 nm) next to the interface is needed for the electrical surface passivation and a C-rich film of adjusted layer thickness serves as anti-reflection coating (ARC).
Si-rich a-Si1-xCx films exhibit good passivation properties on planar (shiny etched) phosphorous diffused n+-surfaces (J0e ≈ 40 fA/cm2), however a strongly deteriorated performance is observed on textured (random pyramids) surfaces (J0e ≈ 400 fA/cm2).
Boron diffused p+-surfaces could not be satisfactorily passivated by a-Si1-xCx. (J0e ≈ 900 fA/cm2). The typical procedure for the extraction of the emitter saturation current density (J0e) from lifetime measurements was found to be unfeasible for the passivation with a-Si1-xCx. This finding is ascribed to the strongly injection dependent surface recombination velocity of a-Si1-xCx at low injection levels independently of the surface doping polarity.
C-rich a-SiyC1-y (y >0.5) layers were optimized regarding their transparency, yield-ing anti-reflection coatyield-ing properties slightly inferior to the ones of typical ARC-SiNx. A low temperature approach for the front side passivation and front contact formation (SE-PassDop) was presented which basically consists of a laser process for opening of the silicon carbide and of subsequent front metallization by Ni/Ag or Ni/Cu plating.
146 Summary
During the laser process, additional dopant atoms incorporated in the front passivation scheme are driven into the surface and allow for the formation of a selective emitter.
First p-type solar cells featuring the PassDop rear and the SE-PassDop front side approach for the passivation and contacting of respective surfaces were fabricated with planar (shiny etched) and textured (random pyramids) front side. A strong parasitic absorption on the front side is mainly attributed to a non-optimized C-rich a-SiyC1-y
layer. Open-circuit voltages of up to 672 mV and 623 mV evidence a high level of surface passivation on planar emitters and a strongly deteriorated performance on textured surfaces, respectively. The feasibility of the metallization concept including the laser process is impressively demonstrated by fill factors of up to 79 %. The incor-poration of a doped, Si-rich a-Si1-xCx film in the emitter passivation scheme proved to be beneficial in terms of series (contact) resistance as well as in terms of front surface recombination.
While the parasitic absorption in the front a-SiCx layer stack should be strongly re-duced by application of an optimized C-rich a-SiyC1-y layer and by further reduction of the thickness of the passivating/dopant containing films, it is not yet clear how to overcome the problems related to the poor electrical passivation of surfaces featuring random pyramids.
As a final, practical remark, passivating a-Si1-xCx is more closely related to amor-phous silicon than to SiNx – from a chemical as well as from an electrical viewpoint.
However, “carbon doping” of the amorphous silicon matrix clearly exhibits a benefi-cial impact on the (low temperature) thermal stability of the film (< 500°C) while maintaining an excellent level of electrical properties.
Deutsche Zusammenfassung
Die in dieser Arbeit untersuchten Siliciumcarbidschichten wurden mittels Plasma unterstützter chemischer Gasphasenabscheidung (PECVD) abgeschieden. Die beiden hierfür verwendeten PECVD Anlagen wurden vorgestellt. Die grundsätzliche Entwick-lung der Schichten wurde an dem AK400M Laborreaktor von Roth&Rau durchgeführt und die entwickelten Prozesse wurden dann auf eine industrielle SINA Inlineanlage derselben Firma übertragen. Für die typischen passivierenden, siliciumreichen a-Si1-xCx Schichten, die in dem AK400M Reaktor unter Benutzung von Hochfrequen-zanregung (HF) abgeschieden wurden, zeigen SIMS Messungen einen relativ geringen Kohlenstoffanteil in der Matrix (< 10 at.%). Dotierstoffdichten in der Schicht von 1×1021 cm-3 für Phosphor und von 5×1020 cm-3 für Bor konnten für typische Gasflüsse von 100 sccm (B2H6 oder PH3 hochverdünnt in H2) nachgewiesen werden. Die a-Si1-xCx Schichten aus der SINA Inlineanlage sind den aus dem Laborreaktor stam-menden Schichten sowohl hinsichtlich ihrer chemischen Zusammensetzungen als auch hinsichtlich ihrer elektrischen Eigenschaften sehr ähnlich.
Eine parallele Untersuchung an a-Si1-xCx/c-Si and a-Si1-xCx/c-Ge Systemen wurde durchgeführt. Aus der Literatur ist bekannt, dass die Hauptunterschiede zwischen der Oberflächenpassivierung von Silicium und Germanium in der Instabilität von (ther-misch gewachsenem) GexOy und in der eingeschränkten Absättigung von Germanium Dangling Bonds durch Wasserstoff bestehen. In Bezug auf die Struktur der abgeschie-denen Schichten zeigen die mittels Transmissions-Elektronenmikroskopie durchge-führten Untersuchungen keinen Einfluss des Substrattyps (c-Si oder c-Ge). Mit zunehmendem C-Anteil werden die siliciumreichen a-Si1-xCx Schichten sowohl optisch als auch in ihrer Struktur weniger dicht, was auf eine erhöhte Anzahl an Mikroporen zurückgeführt wird. Die Untersuchungen zeigen deutlich, dass die Qualität der Ober-flächenpassivierung von c-Si direkt mit der Si-H Bindungsdichte in der Schicht korre-liert ist. Der Beginn des Aufbrechens der Si-H Bindungen bei ungefähr 300°C geht daher einher mit dem Beginn der elektrischen Degradation der Schicht. Dies ist unab-hängig vom Kohlenstoffgehalt und der Dotierstoffdichte in der Matrix. Die Ober-flächenpassivierung von Germanium mit a-Si1-xCx zeigt ein stark abweichendes Ver-halten. Hier ist die thermische Stabilität im Falle von intrinsischen Passivierschichten erheblich erhöht. Die elektrische Degradation findet erst ab Temperaturen von 450°C statt und deutet somit auf eine Entkopplung von Passivierqualität und Wasserstoffge-halt der Schicht hin.
148 Deutsche Zusammenfassung
Die optimale Substrattemperatur während der PECVD Abscheidung wurde zu Topt=270°C ermittelt. Hinsichtlich der Passivierqualität werden hier sowohl auf Sili-cium als auch auf Germanium die besten Ergebnisse erzielt. Ein ähnliches Verhalten für beide Substrattypen zeigt sich bei Abscheidetemperaturen unter oder über dieser Temperatur. Nachträgliches Annealen der Schichten verbessert die Passivierqualität für Tann <Topt und resultiert in einer Degradation für Tann >Topt. Dieses Ergebnis stimmt mit der zentralen Bedeutung von Wasserstoff für die Oberflächenpassivierung von c-Si überein. Eine weitere Schlussfolgerung aus diesen Experimenten kann dahingehend gezogen werden, dass der Kohlenstoff in der a-Si Matrix vermutlich das epitaktische Wachstum verhindert, welches bei der Abscheidung von a-Si bei erhöhten Tempera-turen beobachtet und für die Degradation der elektrischen Eigenschaften der Schicht verantwortlich gemacht wird. Im Falle von c-Ge ist die Qualität der Oberflächenpas-sivierung indirekt mit dem Wasserstoffeinbau bei erhöhten Temperaturen korreliert, da letzterer die Wachstumskinetik der Schicht direkt beeinflusst.
Für die elektrische Degradation der Passivierungsqualität wurden aus isothermen Lebensdauerexperimenten Aktivierungsenergien bestimmt. Für das c-Si/a-Si1-xCx
System beträgt diese Energie ca. 0.5 eV und für das c-Ge/a-Si1-xCx System ca. 2 eV.
Dieses Ergebnis ist ein klares Indiz dafür, dass auf c-Si und c-Ge Substraten unter-schiedliche Passivierungsmechanismen in Bezug auf die a-Si1-xCx Schicht involviert sind. Weiterhin lassen die experimentellen Ergebnisse darauf schließen, dass die ther-mische Stabilität der elektrischen Eigenschaften der Schicht auf ein Gleichgewicht zwischen dem Aufbrechen von Si-H Bindungen und seinem inversen Prozess (Si-H↔Si + H) zurückzuführen sind. Dieses Gleichgewicht wird maßgeblich durch die Verfügbarkeit von freiem atomaren oder molekularen Wasserstoff in der Matrix bes-timmt.
Beidseitig a-Si1-xCx passivierte Siliciumwafer unterschiedlicher Dicke wurden benutzt, um die (“reinen”) injektionsabhängigen Oberflächenrekombinationsgesch-windigkeiten (SRV) zu bestimmen. Das Anfitten der SRV-Daten von n- und p-Typ Siliciumproben unter Verwendung des erweiterten Shockley-Read-Hall (SRH) Models zeigt einen Vorzeichenwechsel der festen Ladungsträgerdichte Qf, wenn der Dotiertyp des Substrates gewechselt wird. Die Existenz von amphoteren Grenzflächenzuständen wurde durch SPV Messungen bestätigt.
Zahlreiche Experimente an dem a-Si1-xCx/c-Ge System deuten auf eine direkte Ab-sättigung der Germanium Dangling Bonds durch Si-/C-Atome hin, was dann eine effektive Unterdrückung der Rekombination an der Germaniumoberfläche zur Folge hat. Die Passivierung von c-Si Oberflächen hingegen ist inhärent mit der Absättigung von Silicium Dangling Bonds durch Wasserstoff verbunden. Obwohl der Einbau von
Deutsche Zusammenfassung 149
Kohlenstoff in die a-Si Matrix eine erhöhte thermische Stabilität der c-Si Passivierung für Temperaturen bis zu 500°C zur Folge hat, konnte im Rahmen dieser Arbeit keine
Kohlenstoff in die a-Si Matrix eine erhöhte thermische Stabilität der c-Si Passivierung für Temperaturen bis zu 500°C zur Folge hat, konnte im Rahmen dieser Arbeit keine