Aloysius F. Hepp, Jeremiah S. McNatt, and John E. Dickman

Glenn Research Center, Cleveland, Ohio

Michael H.-C. Jin

University of Texas at Arlington, Arlington, Texas

Kulbinder K. Banger and Christopher V. Kelly

Ohio Aerospace Institute, Brook Park, Ohio

Angel R. Aquino González and Angus A. Rockett

University of Illinois, Urbana, Illinois

Aerosol-Assisted Chemical Vapor Deposited

Thin Films for Space Photovoltaics

NASA/TM—2006-214445

November 2006

AIAA–2006–4010
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Aerosol-Assisted Chemical Vapor Deposited

Thin Films for Space Photovoltaics

NASA/TM—2006-214445

November 2006

National Aeronautics and

Space Administration

Glenn Research Center

Cleveland, Ohio 44135

Prepared for the

Fourth International Energy Conversion Engineering Conference and Exhibit (IECEC)

sponsored by the American Institute of Aeronautics and Astronautics

San Diego, California, June 26–29, 2006

AIAA–2006–4010

Aloysius F. Hepp, Jeremiah S. McNatt, and John E. Dickman

Glenn Research Center, Cleveland, Ohio

Michael H.-C. Jin

University of Texas at Arlington, Arlington, Texas

Kulbinder K. Banger and Christopher V. Kelly

Ohio Aerospace Institute, Brook Park, Ohio

Angel R. Aquino González and Angus A. Rockett

University of Illinois, Urbana, Illinois
Acknowledgments

Available from

NASA Center for Aerospace Information

7121 Standard Drive

Hanover, MD 21076–1320

National Technical Information Service

5285 Port Royal Road

Springfield, VA 22161

Available electronically at http://gltrs.grc.nasa.gov

Level of Review: This material has been technically reviewed by technical management.

The authors gratefully acknowledge NASA for financial support cooperative agreement NCC3–947 and NCC04AA71A. We also

wish to thank Mr. David Scheiman from the Ohio Aerospace Institute for his assistance with the solar simulator at NASA Glenn

Research Center and Mr. Philip Jenkins also from the Ohio Aerospace Institute for his assistance with the PL measurements. We

are grateful to Dr. Robert Birkmire, Dr. William Shafarman, and the staff at the Institute of Energy Conversion of the University

of Delaware for their help with solar cell fabrication. We also acknowledge the contributions of Dr. Jerry Harris from Northwest

Nazarene University, Dr. William Buhro and Dr. Jennifer Hollingsworth from Washington University, Dr. John Scofield from

Oberlin College, and Mr. Jonathan Cowen from Case Western Reserve University. Finally, we acknowledge

 Dr. Kannan Ramanathan of NREL for providing a shadow mask.
NASA/TM—2006-214445 1

Aerosol-Assisted Chemical Vapor Deposited 

Thin Films for Space Photovoltaics 

Aloysius F. Hepp, Jeremiah S. McNatt, and John E. Dickman 

National Aeronautics and Space Administration 

Glenn Research Center 

Cleveland, Ohio 44135 

Michael H.-C. Jin 

University of Texas at Arlington 

Arlington, Texas 76019 

Kulbinder K. Banger and Christopher V. Kelly 

Ohio Aerospace Institute 

Brook Park, Ohio 44142 

Angel R. Aquino González and Angus A. Rockett 

University of Illinois 

Urbana, Illinois 61801 

Abstract 

Copper indium disulfide thin films were deposited via aerosol-assisted chemical vapor deposition using single 

source precursors. Processing and post-processing parameters were varied in order to modify morphology, 

stoichiometry, crystallography, electrical properties, and optical properties in order to optimize device-quality 

material. Growth at atmospheric pressure in a horizontal hot-wall reactor at 395 °C yielded best device films. 

Placing the susceptor closer to the evaporation zone and flowing a more precursor-rich carrier gas through the 

reactor yielded shinier, smoother, denser-looking films. Growth of (112)-oriented films yielded more Cu-rich films 

with fewer secondary phases than growth of (204)/(220)-oriented films. Post-deposition sulfur-vapor annealing 

enhanced stoichiometry and crystallinity of the films. Photoluminescence studies revealed four major emission 

bands (1.45, 1.43, 1.37, and 1.32 eV) and a broad band associated with deep defects. The highest device efficiency 

for an aerosol-assisted chemical vapor deposited cell was 1.03 percent. 

Nomenclature 

AM0 air mass zero (space solar spectrum) 

AM1.5 air mass one point five (solar spectrum at sea level) 

CIGS copper indium gallium diselenide (CuIn1-xGaxSe2 - a I-III-VI2 material) 

(AA)CVD (aerosol-assisted) chemical vapor deposition 

DAP donor-acceptor pair

EDS energy dispersive spectroscopy

FF fill factor 

Jsc short-circuit current 

η solar cell efficiency 

PL photoluminescence 

SEM scanning electron microscopy 

SSP single-source precursor(s) 

Voc open-circuit voltage 

(GA)XRD (glancing angle) X-ray diffraction 

I. Introduction 

The development of new technologies that will enable the exploration of the universe has been one of the major 

goals of the National Aeronautics and Space Administration (NASA) since its founding in 1958. Technologies that 

will enable the design of new spacecraft, rovers, satellites, and other space vehicles will assist the agency in its quest 
NASA/TM—2006-214445 2

to explore and understand planetary systems, stars, and finally the universe. As part of its efforts to fulfill these 

goals, NASA GRC has been working on the deposition and fabrication of lightweight, thin film solar cells for space 

power (refs. 1 to 9). The ability to deposit thin films on lightweight, flexible polymer substrates will lead to more 

design flexibility, lower launch costs, and extra space and mass for payloads. 

Depositing high quality absorber layers on polymer substrates has been a challenge because polymer substrates 

do not withstand the high deposition temperatures used in standard deposition processes. Temperatures below 

400 °C are usually needed to ensure that the polymer substrate does not suffer thermal degradation. Because of this 

temperature constraint, copper indium disulfide (CuInS2) was deposited using single-source precursors (SSPs) with 

low decomposition temperatures (<250 °C) that were synthesized in-house (refs. 1 to 5). Films were deposited using 

an aerosol-assisted chemical vapor deposition (AACVD) reactor to exploit the lower deposition temperature enabled 

by the simpler decomposition chemistry resulting from the presence of all of the essential elements in the SSPs 

(refs. 6, 7, and 9). AACVD is a simple and inexpensive process that offers the advantage of a uniform, large-area 

deposition, just like metal organic CVD, while also offering the low-temperature solution reservoir typical of spray 

pyrolysis methods. 

Polycrystalline CuInS2 was chosen as the absorber layer because of its high absorption coefficient and its direct 

band gap of 1.5 eV, which is near optimum for AM0 conditions. So far, the highest total area efficiency achieved by 

a CuInS2 cell using co-evaporation techniques has been 12.5 percent (ref. 10). Other alloys of Cu(In,Ga)(S,Se)2, 

such as Cu(In,Ga)Se2 (CIGS), have shown conversion efficiencies as high as 19.2 percent under AM1.5 illumination 

(ref. 11); but the 1.5 eV direct band gap of CuInS2 makes it ideal for solar radiation conversion in space. 

Polycrystalline devices typically outperform single crystal devices created under similar conditions (ref. 12). 

Furthermore, polycrystalline films can be deposited on a broader range of substrates, including polymers, and are 

less sensitive to deposition conditions. CuInS2 thin films are also more resistant to radiation in space than other 

CIGS alloys (refs. 13 and 14); they can be used as a top cell in a tandem structure with CIGS (ref. 15); and they are 

less toxic to process than selenium-containing alloys. 

SSPs have the I-III-VI2 stoichiometry “built in” and offer a clean approach for depositing thin films. One 

advantage is that they offer a number of tunable sites within the complex, allowing various combinations for 

depositing chalcopyrite films of different compositions. The structure of an SSP is shown in figure 1. Early work 

was performed by Nomura (refs. 16 to 19) and Kanatzidis (ref. 20). Buhro and Hepp later demonstrated that 

(PPh3)2Cu(SEt)2In(SEt)2 could be used in a spray CVD process for depositing CuInS2 thin films under 400 °C 

(refs. 21 to 23). Prior to our efforts, very few ternary SSPs were known, characterized, and used in thin film 

deposition; Hepp and Banger have made numerous findings on the synthesis and use of SSPs. A review of this work 

was recently reported (ref. 3). 

Throughout the studies performed at NASA GRC, various processing and post-processing parameters were 

modified in order to determine how these affected morphology, stoichiometry, crystallography, electrical properties, 

and optical properties of deposited thin films. Some of the modifications made included varying the reactor type and 

configuration, deposition temperatures, location of the substrate within the reactor, annealing times, and annealing 

atmospheres. This review will discuss several deposition setups, details of processing experiments, and their impact 

on film properties. A number of characterization techniques including X-ray diffraction (XRD), scanning electron 

microscopy (SEM), and photoluminescence (PL) were used to obtain film properties. Complete working devices 

were also fabricated and tested. 
NASA/TM—2006-214445 3

II. Experimental 

A. Reactor Designs 

Three different reactors were used to deposit CuInS2 films via AACVD; the first reactor design discussed was 

primarily used in our parametric study. This reactor was a horizontal atmospheric pressure hot-wall reactor (reactor 

A, fig. 2(a)) with a plate-type 2.5 MHz ultrasonic nebulizer from Sonaer Ultrasonics. The precursor (1.5 to 3.5 g) 

was dissolved into distilled toluene (50 to 400 ml) and fed into the nebulizer using a syringe pump. This nebulizer 

created an aerosol that was swept into the reactor by an Ar carrier gas at a flow rate of 4 l/min (or SLPM). The 

reactor is a two-zone furnace consisting of a warm precursor evaporation zone and a hot deposition zone. The 

aerosol is converted to precursor vapor and mixes with the reactor gas in the evaporation zone. A solid graphite 

susceptor coated with SiC was placed in the deposition zone. 

The second reactor was a vertical atmospheric pressure, cold-wall reactor (reactor B, fig. 2(b)) with a 

commercial ultrasonic nozzle (Sono-Tek 120 kHz); the third reactor was a horizontal low-pressure hot-wall reactor 

(reactor C) with a pulsed aerosol injection system using a commercial automotive fuel-injector (Ford 2M2EA7B). 

Most depositions were done on 0.87 by 7.6 by 0.1 cm soda lime glass slide substrates (Fisher 12-550 A) coated with 

Mo using a radio-frequency magnetron sputtering system. 
NASA/TM—2006-214445 4

Even though all three reactors share the same precursor delivery system (AACVD), each tool offers specific 

advantages. For example, a cold wall reactor (reactor B) will help prevent decomposition of the precursor before it 

reaches the substrate. A pulsed aerosol injection system at low pressure (reactor C) will allow the film to grow under 

better-defined conditions than in a continuous process because of the minimization of undesirable transient effects 

due to the high volatility of the solvents used (ref. 24). A more detailed description of each of the conditions for film 

growth, including reactor type, precursor type, delivery method, deposition temperatures, growth time, and other 

parameters are summarized in table I. Depositions were done on bare and Mo-coated glass slides from Corning 

(Corning 2947 and 7059). Commercial Mo foil substrates were also used. Post-deposition annealing of film III was 

performed in a tube furnace (Lindberg 54477) at 600 °C for 10 min under N2 flow (4 SLPM). 

TABLE I.—GROWTH PARAMETERS FOR AACVD DEPOSITED CuInS2 FILMS 

Film I II III IV V 

Reactor A B C 

Precursor Solid Liquid 

Precursor concentration (M) 0.01 

Precursor delivery method Carrier gas Vacuum-driven 

Precursor delivery rate (ml/min) ∼1.7 1.5 0.4 

Temperatures (°C) 

Evaporation zone: 

Deposition zone: 

Substrate: 

128 

390 

 390#

120 

360 

 360# 

no zones 

400 

250 

250 

420 

150 

250 

400 

Pressure (torr) atm. ∼12 ∼10 

Ar flow rate (l/min) 4 0.08* 0.15*

Substrate Mo/glass 

(corning 

2947) 

Mo foil Glass 

(corning 

7059) 

Mo/glass 

(corning 2947) 

Substrate size (cm by cm by 0.1 cm) 0.87 by 7.6 2.5 by 2.5 2.5 by 7.6 

Growth time (min) 80 60 250 

Thickness (μm) ∼1 ∼4 ∼1 

Solid precursor: (PPh3)2CuIn(SEt)4

Liquid precursor: {P(n-Bu)3}2CuIn(SEt)4 #

Substrate is located within the deposition zone. *

Direct reading from MKS flow controller (not calibrated). Ar flow is required for the aerosol creation mechanism. 

B. Conditions Used for Atmospheric Pressure Hot-Wall Reactor Parametric Study

The reactor used for this study was the horizontal atmospheric-pressure hot-wall reactor (reactor A). The 

susceptor accommodated three substrates side by side at an angle of 15.5° above horizontal. The leading edge of the 

film experiences a more reactant-rich gas stream and is closer to the bottom of the reactor. The environment of the 

trailing edge, deeper in the deposition zone, is more product-rich and is closer to the axial center of the reactor tube. 

In order to optimize thin film and device quality, a four-phase study was conducted. The temperature of the 

deposition zone, location of the susceptor along the length of the reactor tube, concentration of the precursor 

solution, and post-deposition annealing conditions were varied individually. Initially, the substrate temperature was 

varied between 350 and 425 °C. In the second phase, the measured distance from the end of the reactor tube up to 

the trailing edge of the susceptor were varied from 57.15 mm (2.25 in.) to 127 mm (5 in.). The larger the distance, 

the closer to the evaporation warm zone. For phase three, the precursor concentration was varied between 0.005 and 

0.04 mol/l. The deposition temperatures for the two middle parts of the experiment were maintained at a fixed value 

of 395 °C. Finally, post-deposition annealing of films was performed in the reactor. The evaporation zone 

temperature was set to 120 °C and run-to-run deposition zone temperatures were varied between 450 and 580 °C. 

During annealing, Ar was flowed through the reactor tube at a rate of 10 ml/min. Annealing times varied from 

15 min to 17 hr. The annealing times began when the annealing temperature was reached and ended when the 

cooling process started. It took about 30 min to reach the annealing temperature when started at room temperature 

and a little over 1 hr to cool down to room temperature. For some runs, sulfur vapor was added during annealing by 

placing sulfur (Strem Chemicals, 99+ percent) in a powder-filled crucible in the evaporation zone. 

C. Fabrication of CuInS2 Solar Cells 

Complete CuInS2 solar cells were fabricated at GRC; the cell architecture is shown in figure 3. After coating the 

glass substrates with Mo, CuInS2 films were deposited via AACVD, following procedures optimized for the 
NASA/TM—2006-214445 5

atmospheric-pressure hot-wall reactor (reactor A). Post-deposition, sulfur/Ar-anneals were performed from 450 to 

580 °C for up to 17 hr. The films were then etched in a 1.5 M KCN solution for 1 min at room temperature to 

remove residual Cu2S. Cadmium sulfide (CdS) was then deposited on the films using chemical bath deposition 

(CBD) to form the collecting heterojunction. The CBD used a solution of 0.001 M CdSO4, 1.5 M NH4OH, and 

0.0075 M thiourea (H2NC(S)NH2). The bath was heated to 70 °C. Samples were immersed into the bath for 7 to 

10 min until the solution started turning yellow. Once the CdS had been deposited, films were placed in an 

ultrasonic bath to remove particulates. A layer of ZnO:F was deposited from a 16.5 cm (6.5 in.) diameter F-doped 

ZnO target in a RF sputtering system using pure Ar as the sputter gas. Aluminum (Al) was used as the top contact 

metal. A 0.2 μm thick layer of Al was deposited by thermal evaporation through a shadow mask. The devices were 

isolated by mechanically scribing. The final cell area of the GRC-produced devices was typically ~0.4 cm2

. 

Solar cells were also fabricated at the Institute for Energy Conversion (IEC) following their well-established 

fabrication process (ref. 25). This was done in order to isolate issues related to AACVD deposition of CIS films from 

those associated with the rest of the solar cell fabrication at GRC. For the cells prepared at IEC, only CIS layers 

were deposited at GRC. The rest of the device fabrication, including the coating of the glass substrates with Mo and 

the mechanical scribing, were performed at IEC. Final cell area for IEC-produced devices was typically ~0.45 cm2

. 

Finally, post-fabrication anneals were carried out in air on a hot plate (Cimarec 2) at 150 °C for up to 24 hr. 

D. Characterization Techniques 

Following growth of films and completion of devices, characterization techniques were employed to determine 

the properties and performances of the films and devices. Film thicknesses were measured using a profiler (Sloan 

Dektak IIA or KLA-Tencor HRP 75). Film morphologies were studied using SEM (Hitachi S-3000N and S-800). 

The Hitachi S-3000N includes a built-in EDS (EDAX, Falcon) system used to examine film compositions. 

Compositional measurements were performed in different areas across each cell in order to determine the 

compositional uniformity. The Cu/In ratios were obtained by quantifying the Cu K and In L emission lines using the 

ZAF standardless method. XRD (PANalytical X’pert Pro) was used in order to identify the phases and glancing 

angle X-ray diffraction (GAXRD) was performed using the same instrument in order to identify the phases present at 

the film surfaces. The optical bandgaps of the films were determined by analyzing transmittance measurements 

taken in a UV/VIS spectrophotometer (Perkin-Elmer Lamda-19). Electrical measurements were performed using a 

four-point probe system (Bio-Rad HL5500PC) operated in the van der Pauw configuration and a custom hot-point 

probe system—a commercial soldering iron (Weller TC201) was used as the hot probe. 

Films annealed at 450 °C for 7 hr were characterized using photoluminescence (PL). It has been previously 

shown that annealing can remove structural defects associated with metal ions in films, thereby improving the 

electronic properties of the absorber layer (refs. 26 and 27). CuInS2 films were excited using an Ar ion laser system 

(Coherent Innova 70) with an excitation wavelength of 514 nm. In order to control the laser power density, neutral 

density filters were used. The PL emission detection system consisted of a liquid nitrogen-cooled charge-coupled 

device (CCD) (Horiba Jobin Yvon, Symphony), a germanium detector (EG&G Judson, J16D), and a lock-in system 

(Stanford Research System, SR810DSP). The detectors were connected to a Spex 1269 monochromator, while the 

samples were mounted on a helium cryostat (Janis, STVP-100) in order to lower the temperature to 4 K. Finally, 

devices were tested using a solar simulator at GRC (Spectrolab X-25 Mark II) to obtain current versus voltage 

curves under AM0 illumination. Calibration was performed using a standard GaAs single-junction solar cell. 
NASA/TM—2006-214445 6

III. Results and Discussion 

A. Impact of Reactor Design 

In an initial precursor decomposition study, three different reactors were used to grow CuInS2 films. Hot-wall 

reactors (A (film I) and C (films IV and V)) produced dense, columnar grain growth while the cold-wall reactor 

(film B (III)) yielded porous nanostructures. This can be clearly seen in the SEM micrographs in figure 4, for films I, 

III, IV, and V (see table I for deposition conditions). Figure 4(a), the cross-sectional image of film I, shows a 

columnar grain structure. The columnar grain structure lowers the cross section for photo-excited carrier 

recombination at grain boundaries, thus improving the performance of the solar cells.

Figure 4(b) shows a plan-view image of film III together with an inset of the same region at a higher 

magnification. Under low magnification, only the round structures appeared; however, when the magnification was 

increased smaller nanoparticles (~100 nm) were seen. The nanoparticles may have been created above the hot 

substrate by pyrolysis of nano-droplets generated from the aerosol (ref. 28). Under this model, the ultrasonically 

excited aerosol continues to flow through until it breaks into numerous nano-droplets. Other proposed models 

include nanoparticles coagulating before landing on the growth surface (ref. 29), and liquid deposition taking place 

instead of vapor deposition due to the partial evaporation of the liquid droplet (ref. 30). 

Film IV can be seen in figure 4(c). It exhibits a dendritic microstructure with non-faceted, elongated grains. It 

was previously proven that this was a consequence of diffusion-limited growth (ref. 21). By increasing the flow rate, 

dense and trigonal-shaped faceted grains were obtained for film V (fig. 4(d)). This shape is due to the intersection of 

{211} faces in the chalcopyrite structure. The largest grain sizes obtained using SSPs were of the order of 0.5 μm for 

film V. 

XRD patterns revealed that the films were either (112) or (204/220) oriented. Differentiation between the 

chalcopyrite and sphalerite phases was made by differences in XRD patterns (ref. 31). XRD patterns for all five 

films can be seen in figure 5(a). It was found that films often contained a secondary phase (2θ = 26.5°), believed to 

be indium-rich (In-rich) (ref. 7). The more the film was In-rich, the more likely it was that the secondary phase was 

present with the film being more (204/220)-oriented. EDS measurements (not shown) indicate that the films 

containing the secondary phase were In-rich and that the Cu/In ratio increased when the secondary phase was 

reduced upon annealing (fig. 5(b)). In order to confirm the In-rich nature of the secondary phase, GAXRD was 

performed on the films. The GAXRD pattern (fig. 5(c)) revealed that the secondary phase was concentrated at the 

surface. In order to find out the chemical nature of the phase, films containing the phase were etched in a 10 percent 

aqueous KCN solution for 2 min. It is common to etch Cu-rich CuInS2 films in this solution prior to CdS deposition 

to remove undesired CuSx compounds segregated on the surface during CuInS2 deposition. In addition, it is also 

known that the etch rate of Cu-rich compounds in KCN solutions is much higher than In-rich compounds (ref. 33). 

After etching, GAXRD still showed the presence of the secondary phase on the surface without any change in its 

diffraction intensity, meaning that the phase is not CuSx. The nature of the secondary In-rich phase was also 

confirmed by Raman spectroscopy (ref. 7). 
NASA/TM—2006-214445 7

One way to prevent the formation of the secondary phase is to favor growth conditions of (112)-oriented films 

and avoid the surface kinetic regime where (204/220)-oriented film is favored. In our setup, for any reactor, the 

substrate temperature had to stay around 400 °C in order to promote the growth of (112)-oriented films without a 

high concentration of the In-rich secondary phase. Attempts to grow (204/220)-oriented films without the secondary 

phase were not successful. Another way to reduce this secondary phase is by post-deposition heat treatment. 

Bandgap energies between 1.45 and 1.47 eV were obtained for the films. The band gap energies were estimated 

using plots of (αhν)

2

 versus E (fig. 6), where α is an absorption coefficient estimated from optical transmittance data 

and hν is the photon energy. Figure 6 shows that the band edge sharpens upon post-growth annealing, which is 

ascribed to improved crystallinity and densification (film III). The overall absorption of film V was higher than that 

of film III in the figure, and this could be attributed to the difference in the density between the two films; film V 

had a denser grain structure than film III as shown in figure 4. This could be attributed in part to differences in 

reflectivity; this will be better defined by further work. 

All of the films grown showed p-type conduction regardless of the reactor type, Cu/In ratio, or morphology. It 

was reported that S-rich material shows p-type conduction (refs. 33 and 34); and that the Cu-on-In antisite (CuIn) in 

Cu-rich films is expected to be the major acceptor for p-type conduction (ref. 35). The bulk resistivities ranged from 

0.1 to 30 Ω-cm; Cu-rich films generally had lower resistivity than In-rich films. 
NASA/TM—2006-214445 8

B. Atmospheric Pressure Hot-Wall Reactor Parametric Study 

As discussed above, in this experiment various reactor parameters were manipulated in order to study their 

effects on film growth. In general, the leading edge of the films tended to be smooth, shiny, and dense, while the 

trailing edge tended to be rough, dull, and highly porous. The rough surface morphology can be attributed to the 

stronger diffusion-limited film growth at the trailing edge (ref. 36). By the time the carrier gas (Ar) reaches the 

trailing edge of the substrate, it may be less precursor-rich than the gas flowing over the leading edge – reaction 

occurs on the hot wall of the reactor along the deposition zone. This would make the concentration boundary layer 

thicker at the trailing edge of the substrate, requiring the precursor to diffuse a greater distance to reach the surface 

of the trailing edge and hence the film growth would be more diffusion-limited. SEM images of a rough, dull film 

can be seen below in figure 7(a); the images of a smooth, shiny film can be seen in figure 7(b). 

Run to run deposition zone temperature (350 to 425 °C) was the first parameter varied. Temperature variation 

affected film stoichiometry (fig. 8(a)) and crystalline orientation (fig. 8(b)) while not significantly affecting the 

deposition rate (fig. 8(c)). From figure 8(a), we can see that the films were closest to stoichiometry when deposited 

at 395 °C. Cu-to-In ratios ranged from 0.79 to 0.98, with the highest ratio occurring also at 395 °C. The higher 

deposition zone temperature correlated to increased sulfur content in the films. The crystalline structure was also 

affected by deposition zone temperature. 

CuInS2 films deposited in this study were observed to be crystallographically (220/204)-oriented or (112)-

oriented by XRD. Chalcopyrite solar cells had a higher series resistance for the (220/204)-oriented films (ref. 37). 

From XRD measurements in figure 8(b), we observed that deposition temperatures around 395 °C yielded higher 

ratios of the preferred (112)-oriented films. Finally, deposition rates were not limited by supply of thermal energy 

(temperature variation), therefore film growth was not reaction-limited, see figure 8(c). 
NASA/TM—2006-214445 9

The next parameter varied was the location of the susceptor within the deposition zone. Variations of the 

susceptor location did not affect film stoichiometry, but they did influence the morphologies of the films. When the 

susceptor was moved towards the evaporation zone, denser and smoother, shiny looking films were obtained due to 

a reduced concentration boundary layer (less diffusion-limited) towards the evaporation zone. 

The single-source precursor concentration in the carrier solvent was another of the parameters varied. The 

concentrations of the SSP in toluene were varied between 0.005 to 0.04 mol/l. There were no significant changes in 

elemental composition of the films as a function of concentration, but increasing the precursor concentrations gave 

smoother, shinier films than with lower concentrations. This is due to the fact that at higher precursor 

concentrations, more precursor compound was available to the film, reducing the effects of a diffusion limited 

growth process. 

The final parameter manipulated was post-deposition annealing. It was determined that post-deposition 

annealing improves the elemental composition and crystalline structure of films. For example, after a 5 hr anneal at 

580 °C in a sulfur-rich atmosphere, typical film elemental compositions changed from 51.5 percent S, 24.5 percent 

In, and 24.0 percent Cu, to 52.0 percent S, 23.5 percent In, and 24.5 percent Cu. Annealing increased the Cu-to-In 

ratio which is correlated with improved electrical properties for CuInS2 films (ref. 38). Another advantage of 

annealing in a S-rich atmosphere is that increased S content passivates S vacancies, an undesirable n-type dopant 

(ref. 39). Ordering of the Cu and In atoms within the crystalline structure also increased, producing a better ordered 

chalcopyrite film. CuInS2 films typically occur in either a sphalerite (more disordered) or a chalcopyrite structure. 

Evidence of a more ordered chalcopyrite structure can be seen in figure 9, where the (101) and (211) diffractions at 

17.9° and 37.3°, respectively, can be clearly seen in the annealed film. Another characteristic of the chalcopyrite 

structure not seen in the sphalerite structure is the peak splitting in the (200)/(004) and (116)/(312) diffractions at 

32.2/32.4° and 54.8/55.1°, respectively. The highest efficiency for a cell made with a AACVD-deposited film using 

SSPs was 1.0 percent and the film was annealed under S/Ar at 450 °C for about 7 hr (ref. 8). The highest efficiency 

for a cell made with a non-annealed film was 0.9 percent (ref. 8). Both cells had smooth, shiny, and dense CuInS2

absorber layers. 
NASA/TM—2006-214445 10

C. Post-Deposition Annealing Effects Studied by Photoluminescence 

PL measurements were performed on 76 mm (3 in.) long annealed films grown by the process used in the 

previous experiment. PL spectra of the leading and trailing edges were recorded for as-deposited, S-annealed, and 

Ar-annealed films. As mentioned in previous sections, the leading edges were always more In-rich than the trailing 

edges and, therefore have lower Cu-to-In ratios (0.84 compared to the 1.03). Four different emission bands - PL1 at 

1.45 eV, PL2 at 1.43 eV, PL3 at 1.37 eV, and PL4 at 1.32 eV in figure 10 were identified in both the leading and 

trailing edges of almost all samples (PL1 was not clearly seen in some samples). For the as-deposited films, a near 

band-edge emission can be seen for the trailing edge sample in figure 10. This is probably due to the combination of 

excitonic recombination and other transitions associated with unidentified shallow defect levels (ref. 27). Clear 

differences between the leading and trailing edges were observed from emission bands associated with deeper defect 

levels. For the trailing edges, the intensities of PL3 at 1.37 eV were always higher than those of PL4 at 1.32 eV 

(fig. 10). Finally, the broad, deep-level emission bands spanning from 1.32 to 1.24 eV actually tailed down to 

0.9 eV, as measured by a Ge detector. 
NASA/TM—2006-214445 11

After post-deposition Ar and Ar/S-annealing, the films were studied again using PL. S-anneals reduced the 

relative intensities of the PL1, PL4, and broad emission bands, while Ar-anneals increased the relative intensity of 

the PL1 band. This can be seen in figure 11. We can also see in this figure that S-anneals suppressed the broad near 

band-edge emission from the trailing edge samples. When EDS measurements were performed on the films after Sanneals, an extra 1 to 2 at percent S was found incorporated into the films. On the other hand, when Ar-anneals were 

performed, a 5 to 6 at percent S loss occurred. XRD measurements yielded the same results as above, where more 

ordered chalcopyrite structures were obtained after S-anneals (ref.8). S-anneals yielded higher resistivity films while 

Ar-anneals yielded lower resistivity films. Hot-probe measurements showed that all annealed films showed p-type 

conduction. 

The excitation intensity was also varied in order to determine the effect it had on the PL spectrum. PL1 and PL3 

bands had a blueshift per decade of 3.7 and 5.5 meV, respectively, with an increase in excitation intensity. The 

blueshifts were attributed to donor-acceptor pair (DAP) recombination (refs. 27 and 40). PL2 did not show any 

excitation power dependency, while the analysis of the PL4 band was not attempted because of the uncertainty in its 

precise location. The effect of increasing excitation intensity can be clearly seen in figure 12. 
NASA/TM—2006-214445 12

The blue shift of the bands with increasing excitation intensity is explained as follows. The emission energy for 

the recombination can be given by: 

DA

g A D r

e h r E E E ε ν = − + +

2

( ) ( ) (1) 

where Eg is the band gap, EA (ED) is the activation energy of the acceptor (donor), and e

2

/εrDA is the Coulomb energy 

for the pair at a distance rDA (ref. 40). When the excitation energy increases, the film is flooded with photo-generated 

carriers. As a result, the average separation between the donor and the acceptor decreases causing the blueshift. 

Shallow donors and acceptors for CuInS2 have been previously studied and identified (refs. 39 and 41 to 45). 

Figure 13 includes some of the most energetically favorable intrinsic defects used in our analysis to assign 

transitions to the four different emission bands. Shallow donors include: sulfur vacancy (VS), In-on-Cu antisite 

(InCu), and indium interstitial (Ii). Acceptors include: copper vacancy (VCu), Cu-on-In antisite (CuIn), and copper 
NASA/TM—2006-214445 13

interstitial (Cui). The PL1 emission band at 1.45 eV was assigned to a donor-acceptor transition, because of its 

blueshift upon increasing excitation intensity. VS was attributed to the donor state because of the band’s sensitivity 

to the change of VS concentration upon annealing, and VCu was attributed to the acceptor state because the PL1 

intensity was relatively stronger in leading edge (In-rich) samples. The PL2 band at 1.43 eV was assigned to a 

defect-band transition because there was no noticeable blueshift with increasing excitation energy. The PL3 band at 

1.37 eV was assigned to transitions between metal ions (refs. 39 and 46). Finally, the PL4 band at 1.32 eV was 

assigned to a VS-Cui transition. When the film was S-annealed, this band was suppressed; S incorporates into the 

structure reducing the VS concentration (ref. 46). 

D. Testing of CIS Solar Cells 

The best cell fabricated at GRC was made from a film deposited using a 0.01 M solution. After deposition, a Sanneal at 450 °C for 17 hr was performed. This cell was not annealed after fabrication. The Mo bottom contact layer 

had a sheet resistance of 0.5 Ω/sq and a thickness of 0.7 μm. The CdS layer had an optical transmittance of 

65 percent at a photon energy of 1.5 eV. GRC-fabricated cells used a 1 μm thick, n-type ZnO:F window layer with a 

sheet resistance of 200 Ω/sq and a transmittance of 85 percent at a photon energy of 1.5 eV. Thickness of Al top 

contact was ~200 nm. The characteristic curve for this cell together with cell parameters can be found in figure 14 

and table II, respectively. 

The cell fabricated at IEC is also included in figure 14 and table II. This cell was fabricated from a film 

deposited using a 0.04 M solution and S-annealed at 450 °C for 6 hr and 40 min. The completed cell was then 

annealed under air at 150 °C for about 24 hr. This cell’s window layer, a 70 nm thick i-ZnO layer, had a 

transmittance of over 90 percent at a photon energy of 1.5 eV. The sheet resistance of the double top contact layer 

(2 μm Al/50 nm Ni) was about 1.8 Ω/sq. Both cells had nearly the same efficiency (~1 percent), but the IEC cell had 

a higher shunt resistance, which improved the fill factor (table II). The use of the intrinsic ZnO and the Ni/Al top 

contact improved the quality of the IEC cells by lowering the series resistance and improving the fill factor, as 

previously reported in the literature (ref. 15). It should be also noted that most solar cells prepared with the AACVD

method in the literature had a superstrate structure (ref. 6); there has been no report of a working device with a 

typical chalcopyrite solar cell structure (top contact/window layer/CdS/CuInS2/Mo/ substrate) except for the 

authors’ previous results (ref. 47).
NASA/TM—2006-214445 14

TABLE II.—AM0 OUTPUT CHARACTERISTICS OF SOLAR CELLS FABRICATED 

WITH AACVD-DEPOSITED CUINS2 FILMS 

Fabricator Voc

(mV) 

Isc

(mA/cm2

) Fill factor h 

(percent) 

GRC 309 12.5 0.37 1.0 

IEC 412 7.2 0.45 1.0 

E. Thin Film Deposition of Other Chalcopyrite Materials Using SSPs 

Throughout the course of research at NASA GRC, SSPs for depositing other I-III-VI2 compounds as thin films 

have been prepared (refs. 3 and 4). AgIn5S8 thin films with a thickness of about 1 μm and grain size of 

approximately 0.5 μm were deposited using a SSP which should have afforded bulk AgInS2 (ref. 4). Run to run 

deposition temperatures were varied between 350 and 450 °C but no changes in crystalline structure were observed, 

indicating that the deposition temperature was not a limiting factor. CuGaS2 (112)-oriented films were also 

deposited (ref. 3). These films appeared to be dense with an average grain size of 0.4 μm and bandgap energies of 

2.42 eV. An alloy film of Cu(In:Ga)S2 was also deposited using two analogous Ga and In SSPs in a dual-source 

spray CVD setup (ref. 3). The deposited film’s composition and microstructure varied along the length of the film. 

The atomic percentage of Ga increased as the length increased. Compositions of CuIn0.43Ga0.57S2, CuIn0.38Ga0.62S2, 

and CuIn0.27Ga0.73S2 were obtained for the beginning, middle, and end of the film, respectively. The thermal profiles 

of the two SSPs used for depositing these films were unalike. Using two SSPs with matching thermal profiles may 

help in controlling the composition when depositing multi-ternary films. 

IV. Conclusion 

CuInS2 thin films were grown at low temperatures (<400 °C) via AACVD using SSPs ((PPh3)2Cu(SEt)2In(SEt)2). 

During the growth process, it was noticed that (112)-oriented films were close to stoichiometric with no secondary 

phases present, while (204/220)-oriented films were In-rich and always included an In-rich secondary phase. This 

indium secondary phase was removed by post-processing annealing at 600 °C, thereby increasing the Cu/In ratio. 

CuInS2 films grown always exhibited p-type conduction and bandgap energies around 1.46 eV±0.02. Films with a 

grain size up to 0.5 μm were obtained. 

Variation of the deposition zone temperature affected the film stoichiometry and crystalline structure while not 

significantly affecting the deposition rate. A deposition zone temperature was optimized at 395 °C to produce (112)- 

oriented films without any detectable secondary phases. The susceptor location within the furnace did not affect the 

stoichiometry of deposited films but it did alter morphology. Moving the susceptor towards the evaporation zone 

improved the morphology of the films. Increasing the precursor concentration in the carrier solvent also improved 
NASA/TM—2006-214445 15

the morphology of the films while not altering the stoichiometry. EDS and XRD analyses indicated that after Sannealing the stoichiometry and crystalline structure of the films was enhanced. PL studies revealed four major 

emission bands (1.45, 1.43, 1.37, and 1.32 eV) together with a broad band associated with deep defects. The 

blueshift of the 1.45 and 1.37 eV emission bands with increasing excitation power revealed DAP transitions. The 

broad band together with 1.45 and 1.32 eV bands were reduced upon S-annealing due to passivation of VS. 

These films were then further processed to fabricate working solar cells. The best parameters obtained from the 

cells were the following: the IEC cell had the highest Voc and FF, 412 mV and 0.45, respectively, while the GRC 

cell had the highest Isc with 12.5 mA/cm2

. Devices fabricated at both GRC and IEC showed efficiencies above 

1.0 percent (AM0). The major challenge has been achieving higher Voc, which seems to be a problem for cells 

prepared with AACVD deposited films in general. 

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AIAA–2006–4010

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WBS 953033.01.03.23

23

Aerosol-Assisted Chemical Vapor Deposited Thin Films for Space Photovoltaics

Aloysius F. Hepp, Jeremiah S. McNatt, John E. Dickman, Michael H.-C. Jin,

Kulbinder K. Banger, Christopher V. Kelly, Angel R. Aquino González, and

Angus A. Rockett

Copper compounds; Photoluminescence; Solar cells; Thin films; Vapor deposition;

Pyrolysis

Unclassified -Unlimited

Subject Categories: 23 and 44

Prepared for the Fourth International Energy Conversion Engineering Conference and Exhibit (IECEC) sponsored by

the American Institute of Aeronautics and Astronautics, San Diego, California, June 26–29, 2006. Aloysius F. Hepp,

Jeremiah S. McNatt, and John E. Dickman, NASA Glenn Research Center; Michael H.-C. Jin, University of Texas at

Arlington, 701 South Nedderman Drive, Arlington, Texas 76019; Kulbinder K. Banger and Christopher V. Kelly, Ohio

Aerospace Institute, 22800 Cedar Point Road, Brook Park, Ohio 44142; Angel R. Aquino González and Angus A. Rockett,

University of Illinois, 506 S Wright St., Urbana, Illinois 61801. Responsible person, Aloysius F. Hepp, organization code

RPV, 216–433–3835.

Copper indium disulfide thin films were deposited via aerosol-assisted chemical vapor deposition using single source

precursors. Processing and post-processing parameters were varied in order to modify morphology, stoichiometry,

crystallography, electrical properties, and optical properties in order to optimize device-quality material. Growth at

atmospheric pressure in a horizontal hot-wall reactor at 395 °C yielded best device films. Placing the susceptor closer to

the evaporation zone and flowing a more precursor-rich carrier gas through the reactor yielded shinier, smoother, denserlooking films. Growth of (112)-oriented films yielded more Cu-rich films with fewer secondary phases than growth of

(204)/(220)-oriented films. Post-deposition sulfur-vapor annealing enhanced stoichiometry and crystallinity of the films.

Photoluminescence studies revealed four major emission bands (1.45, 1.43, 1.37, and 1.32 eV) and a broad band associated with deep defects. The highest device efficiency for an aerosol-assisted chemical vapor deposited cell was

1.03 percent.