Category Archives: BACKGROUND

Service life prediction from results of accelerated testing

Accelerated life testing means to quantitatively assess the sensitivity to the various degra­dation factors on the overall deterioration of the performance of the component and its ma­terials.

Figure 6 Change in thermal emittance observed for some reference solar fagade absorber materials during outdoor testing and during accelerated corrosion testing. The corrosivity dose in terms of metallic mass loss of copper at an exposure time is also given for the different tests to illustrate that outdoor performance of those absorbers can be predicted by making use of the equivalent corrosivity dose approach.

Mathematical models are then set up to characterize the different degradation mecha­nisms identified and from the accelerated life test results the parameters of the assumed model for degradation are determined and the service life then estimated.

In Figure 6 is illustrated how the principle of equivalent corrosivity dose in accelerated cor­rosion testing can nicely be adopted in the prediction of the long-term outdoor performance of some solar fagade absorbers. A prerequisite for this is that the accelerated corrosion test correctly simulates the predominating corrosion mechanism occurring under normal outdoor conditions.


The best approach in validating an estimated service life from accelerated testing is to make use the results from the accelerated life tests to predict expected change in material properties or component performance versus service time and then by long-term service tests check whether the predicted change in performance with time is actually observed or not.

The results of validation tests therefore can be used to revise a predicted service life and form the starting point also for improving the component tested with respect to environ­mental resistance, if so required. It should be remembered that the main objective of ac­celerated life testing is to try to identify those failures which may lead to an unacceptable short service life of a component. In terms of service life, the main question is most often, whether it is likely or not, that the service life is above a certain critical value.

In the case studies of Task 27 outdoor tests at different test sites are performed for meas­urement of microclimatic variables and for validating predicted loss in outdoor performance from accelerated test results. Tests are performed by CSTB in Grenoble (France), ENEA in Rome (Italy), INETI in Lisbon (Portugal), ISE in Freiburg (Germany), NRELin Colo — rado/Florida/Arizona (USA), SP in Boras (Sweden), SPF-hSr in Rapperswil (Switzerland) and Vattenfall in Alvkarleby (Sweden). In Figure 7 a view of the test site at INETI in Lisbon is shown.

Figure 7View of the outdoor exposure site with facilities for monitoring of climatic data at INETI in Lisbon


The work in IEA Task 27 on durability assessment of static solar energy materials has shown that it is possible to employ a systematic approach in the evaluation of the expected
service life of the materials studied. Based on the work performed recommended test pro­cedures will be worked out for qualification of new materials with respect to durability.

Figure 7 Results from outdoor exposure of antireflective glazing materials performed at SPF-HSR Rapperswil, Switzerland. The decrease in the solar transmittance with time is due to soiling effects, which vary very much with exposure site.

For recommended durability test procedures to be accepted as international standards, it is of utmost importance to demonstrate their relevance for predicting real in-service long­term performance. We think that the work of Task 27 will meet this requirement.


The authors sincerely want to thank the colleagues and participants in the work of Task 27 on the static solar materials for contributions to this paper: Michael Kohl and Volker Kubler (Fraunhofer ISE Freiburg), Ole Holk (DTU Copenhagen), Gary Jorgensen (NREL, Golden Colorado), Bjorn Karlsson (Vattenfall Utvecklings AB Alvkarleby), Manuel Lopes Prates (INETI Lisbon), Kenneth Moller (SP Boras), Marie Brogren, Arne Roos, Anna Werner (Uni­versity Uppsala), Michele Zinzi (ENEA Rome), and Michele Ghaleb (CSTB Genoble)



Anca DUTA, Transilvania University of Brasov, Center for Sustainable Development, Romania

Marian NANU, Delft University of Technology, the Netherlands

Introduction: The concept of three dimensional(3D) solar cells lately developed, […], represents a promising alternative to silicon solar cell. This type of solar cell is based on a composite p-n heterojunction mixed on a nanometric scale.

Dense and nanoporouse n-type anatase layers are obtained by Spray Pyrolisys. These films solve the electron conduction in a 3D solar cell with the structure TEC8/dense n-TiO2/nanoporous n-TiO2/ p-CuInS2/Au.

Thin layers of TiO2 were obtained using titanium tetra-isoprpoxyde (TTIP) as precursor. The morphology, crystallinity and conductive properties are discussed in conjunction with the deposition parameters.

The photovoltaic response of the 3D solar cell is presented.

The Three — Dimensional Solar Cell

The development of energy-friendly alternatives to the silicon solar cells, efficient and with short payback time, represents a research topic formulated in the past 15 years. The solar cells involving only solid state compounds, not silicon based, are now-a-days developed: the ETA cells, [2] (where a thin nanoporous n-type and p-type semiconductors) represents an advent of the dye-sensitized (Gratzel type), [3] solar cell.

With the advent of polymer bulk heterojunction and ETA solar cells, a three dimensional (3D) solar cell represents a further development where the n-type nanoporous wide band gap semiconductor thin layer is infiltrated with a p-type, light absorbing nanoporous semiconductor. The large interpenetration nanoporous heterojunction overcomes the drawback of the low conduction in solid state.

In a 3D solar cell the n-type semiconductor is usually anatase TiO2, chosen because its inert chemical behavior in various environment. This condition is necessary to be fulfilled since the p-type semiconductor, CuInS2 (with the band gap of 1.55 eV) is obtained in deposition and reductive annealing conditions that can affect a more reactive oxide (e. g. ZnO considered state-of-art in the ETA cell). For avoiding shunts at the back contact interface, the anatase consists of two layers: a dense thin layer, with low flexibility (100 nm) and the nanoporous matrix (1000 nm) able to be infiltrated with CuInS2.

Literature reports different techniques for obtaining solid state solar cells: The ETA cells with the highest efficiency (2%) are obtained by dipping of microporous anatase TiO2 in liquid precursors of Cd and Te, [4]; CuInS2 for solar application was obtained by vaporization of the metal precursors followed by annealing. Recently, CVD and AlD deposition was reported for CuInS2 in a 3D solar cell, [1].

Spray Pyrolisys Deposition (SPD) is a simple and attractive technique and the aim of developing a 3d solar cell using only this procedure is already formulated.

This paper investigates the possibility of obtaining dense and nanoporous anatase TiO2 layers using SPD. The deposition parameters are discussed in conjunction with the morphology, structure and conduction properties of the obtained layers.


Absolute ethanol (EtOH, 99.99% Aldrich) solutions of TTIP (97% Aldrich) were prepared and acetylacetonate (AcAc 99+%, Aldrich) was added in order to obtain regular morphologies.

Dense thin TiO2 films were prepared by spraying, at 350oC, a mixture with a volume ratio TTIP : AcAc : EtOH = 1 : 1.5 : 22.5.

The nanoporous films were prepared by SPD at 450oC, using a mixture with the volume ratio: TTIP : AcAc : EtOH = 1.5 : 1 : 18.33. The substrate was conductive TCO/TEC8 glass (SnO2:F) with a low internal resistance but a rough surface morphology.

A complete ETA cell was developed having a dense n-type anatase layer deposited by SPD, a buffer layer of In2O3 and the absorber — p type CuInS2 both layers being deposited using Atomic Layer Deposition (ALD) as described elsewhere, [5].

The films were characterised as such and after annealing in reductive atmosphere for six hours (hydrogen gas, at 450oC and 1 mbar).

The crystalline structure was investigated by XRD (Bruker D8, CuKal ).

The morphology of the platinum sputtered layers was investigated by Scanning Electron Microscopy (Joel JM 5800 LV). Luminescence spectra were recorded with a home-built set-up (TU Delft) using a Spectra Physics Millennia Nd:YVO4 laser with a wavelength of 532 nm. The recordings were done in a backscattering mode, using a set of notch filters to remove the Rayleigh scattering and a liquid nitrogen flow cooled the CCD camera (Princeton Instruments LN/CCD-1100PB). A Spex 340E monochromator equipped with a 100 grooves/mm grating was used. Corrections for the filters, the sensitivity of the CDD camera and the monochromator are applied. The same set-up, used in a backscattering mode using a set of notch filters and a Spex 340 E monochromator equipped with a 1800 grooves/mm grating, was used for the Raman spectra.

A Solartron 1286 Electrochemical Interface was used for potentiostatic control and to conduct the Mott Schottky and flat band measurements.

For the ETA cell, the current-voltage (I-V) curves are recorded with a DC source Meter (Keithley, Model 2400) in tha dark and under illumination. A calibrated solar simulator, SolarConstant 1200 (K. H. Steuernagel Lichtechnik GmbH) is used as artificial light source.


Cyclic voltammograms were taken on the various films. In general, the shapes of the voltammograms changed depending on the specific additive, but the main features characteristic of the pure oxide tended to prevail. The charge capacity—and hence the magnitude of the electrochromism—is influenced by the potentiodynamic range, particularly the magnitude of the voltage for full coloration, Ucol. It appeared that similar charge capacities (from 15 to 20 mC/cm2) could be obtained provided that Ucol was varied by 0.05 to 0.1 V when additives were present. This shift is insignificant for electrochromic device applications.

The electrochromism of the films was characterized in terms of a coloration efficiency CE obtained from voltammograms and optical data according to [14]

(1- Rc )2 T

(1- R )2 T,


where ДО is the total charge that is exchanged and subscripts b and c refer to bleached and colored states, respectively. The total exchanged charge was

measured for the cathodic bleaching rather than the anodic coloration in order to minimize the effect of oxygen evolution.

Figure 1(a) shows spectral CEs for a nickel-oxide and a nickel-aluminum-oxide film. Both of these films were optimized by an additional introduction of some hydrogen during the deposition. It is noteworthy that the films have CEs that are much larger than those reported in the literature [1]. It should be emphasized that the CEs in nickel oxide and nickel-based oxides are intimately related to the conditions under which the depositions take place, which implies that the results of the enhanced CEs are related to features such as crystallinity, grain size, porosity, and contents of oxygen and hydrogen in the films. The film with the largest effective grain size presents the lowest CE, thus supporting the idea that a large inner surface area of the film is connected with the electrochromic activity, and that the coloration process takes place in the outermost parts of the grains. This fact—together with good crystallinity and high porosity, along with optimized quantities of oxygen and

(a) (b)

Wavelength (nm) Wavelength (nm)

Figure 1. Spectral coloration efficiency (CE) of nickel-oxide-based (a) and iridium-oxide — based films (b) of the shown compositions. NiAlO and IrTaO indicate that Al and Ta are present in the oxides but do not specify the amount.

hydrogen—gives highly efficient electrochromic films [15].

Figure 1(b) shows spectral CEs for two different iridium-oxide-based films. Clearly iridium-tantalum oxide has a lower CE than pure iridium oxide for some wavelengths. However it does not follow that Ir oxide is superior to IrTa oxide in applications, since a more crucial property may be the bleached state transmittance, as we elaborate below.

Optimized TI structure for solar wall applications

Fig. 1 shows a sketch and a fotograph of optimized small-celled TI structures based on polymer films. The structures, which consist of a continuously produced small-celled lamella based on a flat film joined to a corrugated film are either stapled or manufactured in the form of rolled structures. Due to the use of high-quality films structural defects on the surface and within the TI structure were avoided. Based on investigations of more than 20 different polymer film types [2, 3], cellulose acetate films (CA and CTA) were identified as outstanding polymer materials for TI wall applications. Using 30 pm thick cellulose acetate films small-celled lamellae with widths between 100 and 135 mm and a heigth of 5 mm were produced. The material fraction of the lamellae and the stapled or rolled structures thereof were about 1.5 v%. For a 100 mm thick stapled structure a hemispherical solar transmittance of 0.80 and a heat conductance 0.85 W/(m2K) were measured. For 135 mm thick structures, which were used for the application demonstration object, an hemispherical solar transmittance of 0.74 and an heat conductance of 0.74 W/(m2K) were calculated using the program GWERT [5] and polymer film properties as input data [2,3]. The remarkable performance property profile of cellulose based materials for TI applications with maximum service temperatures of about 100°C is related to the high solar transmittance of CA films in combination with a high infrared absorptance due to a high density of functional carbon-oxygene single bonds within the molecular structure of CA. However, this high density of carbon oxygene functional groups results in a high moisture uptake of cellulose based materials. For the fully substituted cellulose triacetate (CTA) the maximum moisture uptake is about 3 m%.

Fig. 1. Stapled and rolled small-celled structures based on polymer films

EXPERIMENT. Thin film deposition and photoelectron spectroscopy

The titanium oxide, silicon oxide and aluminium oxide thin films were prepared by reactive magnetron sputtering in a high vacuum deposition chamber using an Ar/02 gas mixture. The mass flow ratio is set to 35:5 for titanium and silicon oxide and to 37:3 for aluminium oxide. The magnetrons are driven by bipolar-pulsed power for the Ti target (50 kHz at 200 W) and for the Al target (50 kHz at 150 W) and by medium frequency RF power (13.5 kHz at 100 W) for Si target. During the thin film deposition carried out at room temperature, the grounded substrates face the target at a distance between 5 and 8 centimeters. A working pressure of around 3 x 10-3 mbar is adjusted by throttling the pumping system. Ti02, Si02 and Al203 films are deposited on glass AF45 and monocrystalline silicon (with its native oxide) substrates 4×4 cm2 for the characterization optical techniques. For the in-situ photoelectron spectroscopy thin films are deposited on sputter cleaned copper substrates. The high vacuum deposition chamber is connected to an ultrahigh vacuum (UHV) electron spectrometer. Samples can be transferred from one system to the other without breaking the vacuum to get chemical information about the deposited films. The electron spectrometer is equipped with a hemispherical analyzer (Leybold EA 10/100) and a X-ray source for core level spectroscopy (X-ray photoelectron spectroscopy XPS: Mg Ka excitation, ho = 1253.6 eV). The typical resolution is 0.8 eV for the XPS measurements. A gold sample with the Au 4f7/2 core level signal at 83.9 eV binding energy is used as a reference for the electron energy calibration.

Ti 2p, Si 2p, Al 2p and 01s core levels were recorded in the case of Ti02, Si02and Al203 to determine the chemical composition for each layer. Atomic concentration ratios were calculated by integrating the peaks area after subtracting a Shirley background [12]. From the photoionization cross-sections given by [13], the atomic concentration at the films surface is calculated using UNIFIT program [14]. In our deposition conditions, the stoichiometry ofthe Ti02, Si02and Al203 films is reproducible.

Photoelectrochemical Properties

The band gap was determined by absorption [3] and photoelectrochemical [7] methods and was found to be 1.7 and 1.75 eV respectively. The quantum yield vs. stimulation is shown in Fig. 1. A further reference [8] gives 1.5 eV band gap and indirect band transition. Ref. [9] gives 1.9 eV for band gap from photoluminescence measurement at 10 K. These strongly different photoelectrochemical and photoluminescence results are refined in this work.

The impedance measurements were performed in an electrochemical cell under potentiostatic control. The electrolytes were 0.05 M H2SO4 and 1 M KOH solutions. The impedance analysis was carried out with the perturbation of some mV. The modelling of an electrolyte — semiconductor junction is a difficult problem because the values of the circuit elements exhibited frequency dependence. In this work we determined the proper values of equalent circuit components with their physical meaning for the transfer function of the junction. The parameters of the equivalent circuit are very importat to know for device applications. A simple equivalent circuit with physical meaning was appropropriate for the calculation [5] which is contain three paralell branches. One of this branches is a resistance R|. The second branch is a swinging circuit R2C2 and the third branch is a capacity C3.

[oh»: The evaluation of the measured data

was carried out with the help of a computer program developed by us in Turbo Pascal language. The transfer function of the equivalent circuit has three solutions (one zero and two poluses). In the first step these three roots weree fitted in the same time with the help of least square method. The minimum of the error is determined with the help of gradient method, until the error become less than 1 Hz. The value of the constant in the transfer function was determined from the amplitude diagram with similar method [10]. The R1 resistance represents the charge transfer that is the electrochemical reaction at the interface. The value ranges between 6.2 and 7.1 kQ /mm2. R2 and C2 represent the surface levels and deep centres where the values

range between 0.6 and 1.6 kQ/mm2 and 5.3 2

Fig. 2. The amplitude (i) and phase (ii) diagrams of the Cd4GeSe6 and 0.05 M H2SO4 junction where the parameter of the model network are R1 = 51,100 Q (7229 Q/mm2), R2 = 4,227 Q (598 Q /mm2), C2 = 53.8 nF (7.61 nF/mm2), C3 = 4.4 nF (0.62 nF/mm2).


This work was supported by Hungarian National Scientific Foundation (OTKA) through Grant No. T 037509, which is very acknowledged.

New Features in TRNSYS16-New multi-DLL architecture

In TRNSYS 15, adding a new component (Type) required to rebuild the Fortran DLL that included all the TRNSYS Types and the kernel. In version 16, the capability of TRNSYS to adapt to new types of simulation problems has been pushed beyond new limits: Using the Dynamic Link Library (DLL) technology, it is now possible to easily add a new component, written in any programming language, as a Windows™ DLL. This way, component models created by different teams can be used together without even using a compiler: copying the DLL provided by the model author onto the hard disk is sufficient to run it!

TRNSYS 16 is split in several DLL’s that are loaded when TRNSYS is launched. TRNSYS then searches through all the DLL’s for the components that are used in the simulation. This process allows to easily add a new component to TRNSYS by dropping a pre-compiled DLL into the right directory.

The principle of the multi-DLL project is illustrated in the Figure here below.

The TRNSYS kernel only loads the components that are used in a simulation and manages conflicts between DLL’s (e. g. Types defined in more than one DLL). This makes the DLL much smaller in most cases.

It is still possible to recompile TRNSYS 16 as a single DLL for advanced users who give more importance to easy debugging.

Optical, Structural and Kinetic Properties of a New Photoelectrochromic Device

Anneke Georg, Freiburg Materials Research Centre, Freiburg, Germany Andreas Georg, Fraunhofer Institute for Solar Energy Systems, Freiburg, Germany Ursa Opara Krasovec, Faculty of Electrical Engineering, University of Ljubljana, Ljubljana, Slovenia

Photoelectrochromic windows represent a special kind of switching windows. The energy for colouring is provided by sunlight, so that a voltage supply is not required. The transmittance can be decreased on illumination and can be increased again in the dark. In contrast to photochromic devices, the system is externally switchable under illumination.

Our photoelectrochromic window consists of several components: a dye-covered nanoporous TiO2 layer, which is situated on a nanoporous electrochromic layer, such as WO3, two glass substrates coated with a transparent conductive oxide, of which one is coated with Pt, an iodide/tri-iodide redox couple and Li+ ions in a solid ion conductor. All the layers can be kept quite thin, so that they are transparent. The pores of the TiO2 and WO3 layers are filled with the electrolyte.

This configuration is a particularly advantageous combination of the dye solar cell and an electrochromic element. The colouring time is independent of the area, the transmittance can be varied also in the illuminated state, and the system can also be switched by an auxiliary external voltage. Initial samples with solid electrolyte change their visible transmittance from 62 % to 1.6 %, their solar transmittance from 41 % to

0. 8 %. The time for colouring and bleaching is about 10 minutes.


Photoelectrochromic systems combine electrochromic layers [1, 2] and dye solar cells [3, 4]. Electrochromic layers change their transmittance reversibly when electrons and cations are injected. In photoelectrochromic systems, the dye solar cell provides the energy for the coloration of the electrochromic layer. Thus, the transmittance of the photoelectrochromic device can be decreased under illumination and can be increased again when illuminated or in the dark. An external voltage supply is not required. Applications of these devices include, for example, switchable sunroofs in cars or smart windows in buildings.

We developed the photoelectrochromic configuration illustrated in fig. 1, which is a particularly advantageous device. It consists of several components (fig.1): a dye-covered nanoporous TiO2 layer, a porous electrochromic layer, such as WO3, two glass substrates coated with a transparent conductive oxide (TCO), of which one is coated with Pt, an iodide/triiodide redox couple and Li+ ions in an organic solvent. Both the TiO2 and the Pt layers can be kept quite thin, so that they are transparent. The pores of the TiO2 and WO3 layers are filled with the electrolyte.


During illumination (upper part of fig.1), a dye molecule absorbs a photon of the incident light. Then an electron is rapidly injected from the excited state of the dye into the conduction band of the TiO2 and diffuses to the WO3. Ionised dye molecules are reduced by I" in the electrolyte according to the reaction: 3I" ^ I3" + 2e". Li+ ions intercalate into the WO3 and keep the charges balanced. Because of the injection of electrons, the WO3 changes its colour from transparent to blue.

If electrons are allowed to flow via an external circuit from the WO3 via a TCO layer to the Pt electrode (lower part of fig.1, external switch closed), then the Pt catalyses the reverse reaction, i. e. the reduction of I3" to I". Li+ leaves the WO3, and the WO3 is bleached fast. This process occurs also during illumination. If the external switch is open, electrons can leave the WO3 only by loss reactions. This process is very slow.

With a liquid electrolyte, the device’s visible (solar) transmittance under 1000W/m2 of illumination changes from 51% to 5% (35% to 1.5%) with switching times of about 3 minutes. Using a solid electrolyte, a visible transmittance change from 62% to 1.6% and a solar transmittance change from 41% to 0.8% are achieved with switching times of about 10 min. The colouring time is independent of the area.

An alternative photoelectrochromic configuration was first published in [5]. The colouring and the bleaching are competing processes, because the bleaching is possible only via loss reactions. Therefore, either fast colouring and bleaching with a small transmittance change [5] or a large transmittance change with slow bleaching is achievable [6], or an external voltage is used for bleaching [7]. In our new device, the materials can be optimised for colouring and bleaching independently, so it simultaneously allows fast colouring and bleaching, and high contrast [8].

In [8] we introduced this new device and discussed the differences to the alternative photoelectrochromic system and the advantages of our new system.

Experiments with different layer configurations of photoelectrochromic devices were reported in [9]. From these experiments we concluded that the loss reactions of electrons from the TiO2 can be neglected compared to the loss reactions of electrons from the WO3.

We investigated both liquid electrolytes [8,9] and solid electrolytes [10]. Liquid electrolytes allow a faster switching, but need good sealing to be stable on the long term, whereas

solid electrolytes, especially polymer electrolytes, show slower switching properties but are more suitable for most window applications.

Results and discussions

The morphology of the layers showed, both in dense and in nanoporous layers the nanosize of the grains, with a different compaction degree, Fig. 1a and b.

Fig. 1b Nanoporouse anatase matrix

The TiO2 anatase structure was identified in the XRD spectra, Fig. 2. The thin dense layer (100 nm), developed at a lower temperature, shows a peak corresponding to Ti8Oi5, a compound involving also Ti3+. The nanoporous thicker layers (1000 nm) consist only of anatase.

The efficiency of the final solar cell depends on the good conduction along each layer but limitative is the back contact therefore the structure of the dense layer must be free of pinholes for avoiding shunts. The I-V curves, Fig. 3, confirmed that the layers are electron conductive and that thin dense films without pinholes can be obtained.

Fig. 2. XRD spectra of the dense and nanoporous anatase TiO2 layers (* corresponding peak to ТІ8О15)

Fig. 3 Current — Voltage (I-V) curve of the dense anatase layer

Experiments proved that layers thinner than 100 nm, deposited on TEC8, exhibit discontinuities and have a poor conduction.

The defect concentration was calculated in the dense films using the Mott Schotky, Fig. 4. Using the simplified formula (1) the donor density in the space charge region is calculated from the slope of the curve Csc"2 as a function of electrode potential (V), representing the Mott Schottky plot:

1 _ 2 j_

Nd A2eeQ6r H

where H is now the slope of the Mott-Schottky plot and the values used for the constants in eq. (1) are : e=1,610-19C; eo=8.8542i0-12 Fm-1; er=55 and A=3.1410-4 m2.

Fig. 5 The impedance spectrum of the dense TiO2 Fig. 6 The equivalent circuit of the

dense TiO2

The impedance spectrum of the dense anatase film, annealed in oxygen (air), Fig. 5, revealed the existence of the shallow defects (high frequency measurements 1 MHz) while the low frequency signals (10 kHz) are the consequence of the deep defects. The RC equivalent circuit will be modified accordingly, Fig. 6.

The subsequent layers in the 3D cell are deposited at high temperature in an oxygen-free atmosphere. In order to test the behaviour of the anatase layers in the reductive environment, annealing in H2 was investigated.

Annealing in hydrogen increases the oxygen vacancies concentration. Annealing in oxygen exhibits the lowest oxygen vacancies concentration and according to the calculations based on the Mott Schottky plot, the value of the donor density is of 1016/cm3. The Raman spectra of both types of layers did not exhibit any change after hydrogen treatments confirming that the n-type anatse structure is preserved, Fig. 7.

The defect concentration is, however, sensitive to hydrogen annealing, as the photoluminescence measurements show, Fig. 8. The broad peak corresponding to the dense or nanoporous layers may be correlated with the defects existent in the anatase structure as grown via SPD: oxygen vacancies and interstitial titanium ions (+3 and +4). The sharper peaks and the shift of the maximum of the PL spectra to lower values after annealing will than be correlated with the modification of the defects ratio mainly due to the modification in the oxygen vacancies concentration. The XRD measurements showed no increase of the Ti8O15 peak after annealing confirming that interstitial titanium is practically not reduced in the hydrogen atmosphere, in the working conditions.

The photovoltaic response of a 3D cell with the structure: TEC8/dense n-type anatase/ nanoporous n-type anatase/ Al2O3 + In2S3 buffer / p — type nanoporous CuInS2/Au was recently reported, Fig. 9, [1]. The energy conversion efficiency of the cell with a geometrical area of 3.14 x 10-2 cm2 is about 4%. Under AM 1.5 irradiation from a calibrated source the cell has the open circuit voltage VOC of 0.49V, a short-circuit current, ISC, of 18mA/cm2 and a fill factor of 0.44.

Fig. 8

of TiO2 layers,

T= 10K


Using Spray Pyrolisys Deposition the porosity of nanostructurated n-type anatase thin layers can be controlled. By modifying the temperature and precursors composition dense and nanoporous layers can be obtained; the layers are not chemically reactive at high temperature, in hydrogen atmosphere. The layers are electron conductive and can be consequently used for developing 3D solid state solar cells using a complete aerosol technique.


5. M. Nanu, J. Schoonman, A. Gossens, Adv. Mat., 2004, 16, 5, 453

2. K. Tennakone, G. R.R. A. Kumara, I. R.M. Kottegoda, V. P.S. Perera, G. M.L. Aponsu, J. Phys. D, 1998, 31, 2326

3. B. O’Reagan, M. Gratzel, Nature, 1991, 353, 737

4. K. Ernst, A. Belaidi, R. Koenenkampf, Semicond. Sci. Technol. 2003, 18, 475

6. M. Nanu, L. Reijnen, B. Meester, J. Schoonman, A. Goossens, Chem. Vap. Deposition, in press.

7. A. Duta, M. Nanu, I. Visa, Sol. Energ. Mat. Sol. Cells, submitted for publication

Minimized Absorptance

Figure 2 shows spectral absorptance A(A) calculated from

A(A) = 1 — T(A) — R(A).

The data in panel (a) represent the bleached state for the nickel-based oxides mentioned before. Prior to the measurements, the films were cycled ten times in a 1 M KOH solution to stabilize the properties. A significant decrease of A(A) was found at short wavelengths for additives being Mg, Al, Si, Zr, Nb, and Ta, whereas the films containing V and Ag did not show any improvement in their optical properties compared to those of pure nickel oxide.

(a) (b)

The strong absorptance at A < 350 nm is due to the semiconductor band-gap, which appears to be widened as a consequence of the addition of Mg, Al, Si, Zr, Nb, or Ta. On the other hand, the addition of V or Ag narrowed the band-gap. Some weak absorption features can be discerned in the spectral data; they are possibly associated with crystal-field effects [16,17]. Alternatively, the additives may affect
optical absorption caused by defects such as vacancies, over-stoichiometry, grain boundaries, etc. Thermodynamically stable nickel oxide is a p-type conductor due to excess oxygen [18,19]. It is then plausible that the p-type conductivity and the residual optical absorption in the bleached state originate from the same electron states, and this may explain why films of pure Ni oxide cannot be made completely colorless. When Al is added, for example, it can act as a donor of electrons and fill the electron (hole) states on Ni, thereby reducing the residual absorption. The addition of V, on the other hand, may provide acceptor states whose effect would be to enhance the residual absorption.

Figure 2(b) reports comparative data for iridium-based films, cycled in 1 M propionic acid. It is found that additives of Mg, Al, Ta, and Zr tend to lower the absorption, as may be understood by the same arguments as those applicable to the nickel-based oxides.