QUARTERLY TECHNICAL REPORT

28 April 2005 to 27 July 2005

NREL Subcontract ZXL-5-44205-11

Principal Investigator: J. David Cohen

Organization: University of Oregon

We report results of our research activities under NREL Subcontract
ZXL-5-44205‑11 during the first quarter of Phase I. During this quarter
we concentrated our work on two projects primarily related to the issues
of Narrow Gap Materials and, to a lesser extending, metastable
degradation. This work continues our ongoing studies of the HW a-Si,Ge:H
alloys produced at NREL, and the properties of nanocrystalline samples
(nc‑Si:H) obtained from United Solar Ovonics Corporation.

At the end of our previous NREL Subcontract we had reported results on a
series of HWCVD a‑Si,Ge:H deposited at NREL using a filament temperature
maintained at a lower filament temperature than those produced in
previous years. These samples exhibited much improved electronic
properties; in particular, Urbach energies at or below 45meV, even for
samples with Ge fractions approaching 50%. Defect defect densities were
correspondingly low, with values in the mid 10^15^ cm^-3^ range for Tauc
gaps near 1.5eV. These properties represented a dramatic improvement
over any HWCVD a‑Si,Ge:H reported previously and, because the growth
rates exceeded 1.5Å/s for these samples, seemed to offer a potential
solution to the problem of increasing the growth rates of the a‑Si,Ge:H
alloys while maintaining good electronic properties.

In spite of the strong evidence that the properties of these HWCVD alloy
materials were so good, initial attempts by NREL with the help of United
Solar to demonstrate good cell performance with these materials were
unsuccessful. A possible explanation was suggested by some SIMS analysis
at NREL showing that the more recently deposited a‑Si,Ge:H material
(including that incorporated in the cell fabrication studies) had a
nearly 10-fold increased level of oxygen contamination. At the beginning
of the current Subcontract period we carried out a comparison to find
out whether the higher oxygen films did indeed exhibit poorer electronic
properties.

Results from transient photocapacitance (TPC) spectroscopy for a pair of
29at.% Ge samples are shown in Fig. 1, and for a pair of 47at.% Ge
samples are shown in Fig. 2. We clearly observe much broader bandtails
(larger Urbach energy) for the alloy samples with higher oxygen levels,
and also a larger shoulder at lower optical energies indicating a higher
concentration of deep defects. In addition, drive-level capacitance
profiling (DLCP) measurements, while indicating a deep defect density
for the lower oxygen 29at.% Ge sample of 4.0±2.0 × 10^15^ cm^-3^,
indicated nearly an order of magnitude larger deep defect density for
the higher oxygen 29at.% sample, of roughly 2.6±0.25 × 10^16^ cm^-3^.
However, this sensitivity to oxygen impurities is somewhat surprising
given that similar levels of oxygen have not been found to affect the
properties of a‑Si:H at all dramatically. \[[^1]\]

The performance of solar cells depends crucially on the efficiency of
the minority carrier collection and we are able to estimate this by
comparing the *ratio* of the TPC and the transient *photocurrent* (TPI)
spectra in the bandtail regime. To do this we align the spectra at the
lowest optical energies (in the deep defect band region) where only
majority carrier optical transitions are expected to take place. Then
the ratio in the bandtail region, R, where the optical transitions are
expected to generate equal numbers of electrons and holes, should be
equal to (1 - η)/(1 + η), where η = n~h~/n~e~ is the minority carrier
collection fraction, while n~e~,and n~h~ are the respective numbers of
photo-generated electron and holes leaving the depletion region.\[\]
Previous studies that examined this ratio in the best PECVD grown
a-Si,Ge:H materials returned values of at least 10 or more \[9\].

  ---------------------------------------------------- ----------------------------------------------------
  ![](media/image1.wmf){width="3.2256944444444446in"   ![](media/image2.wmf){width="3.2256944444444446in"
  height="4.013194444444444in"}                        height="4.00625in"}

  **FIG. 1.** Comparison of TPC Spectra for two 29at.% **FIG. 2.** Comparison of TPC Spectra for two 47at.%
  Ge HWCVD grown alloy samples with lower and higher   Ge HWCVD grown a‑Si,Ge:H alloy samples with lower
  oxygen levels. In the latter case the Urbach energy  and higher oxygen levels, respectively. We again
  is distinctly larger, and the deep defect band is    note the broader bandtail for the oxygen
  significantly more pronounced.                       contaminated sample.
  ---------------------------------------------------- ----------------------------------------------------

In our previous report we had attempted such a comparison for the lower
oxygen 29at.% Ge sample, but had found very little evidence for hole
collection. However, we later suspected that, to enable reliable
comparisons of the TPC and TPI spectra, we needed to examine films
deposited onto a p^+^ c-Si substrate instead of the standard stainless
steel substrates. This allows our measurements to be performed on the
buried barrier junction at the substrate interface, rather than a top
Schottky barrier contact. The latter type of barrier contact, although
suitable for the TPC measurements, can yield inaccurate TPI spectra for
two reasons: (1) a large additional current signal arises from internal
photoemission near a Schottky contact, and (2) defect states near the
Schottky barrier interface can give a large contribution to majority
carrier current signal (even though they have little effect on the
capacitance signal) due to the different spatial sensitivities of the
two measurements. Indeed, we had determine from our DLCP measurements
that the defect densities were not constant throughout the film for
these HWCVD grown samples and, instead, had much larger deep defect
densities near the top surface of the film (probably due to the increase
in deposition temperature as growth progressed). Such anomalously high
defect densities would then tend to obscure the differences between the
TPC and TPI spectra due to the minority carrier collection.

  -------------------------- ---------------------------------------------------
  **FIG. 3.** Comparison of  ![](media/image3.wmf){width="3.479861111111111in"
  360K TPC & TPI spectra for height="4.521527777777778in"}
  a 29at.% Ge alloy sample   
  deposited by HWCVD onto a  
  p^+^ c-Si substrate. The   
  junction between p^+^ c-Si 
  substrate and intrinsic    
  a‑Si,Ge:H is reversed      
  biased for both the        
  measurements. The          
  difference between the two 
  types of spectra in the    
  bandtail region is related 
  to the fraction of         
  collected photoexcited     
  holes compared to          
  electrons. In this case    
  the ratio of the two       
  signals in the bandtail    
  region (of roughly a       
  factor of 7) indicates a   
  relative hole collection   
  efficiency of 75%.         

  -------------------------- ---------------------------------------------------

Both types of spurious effects could be minimized at the p^+^ c-Si
junction by a special cleaning procedure of these substrates. Such
substrates were prepared at Oregon and then sent to NREL for a‑Si,Ge:H
film deposition. The resultant TPC and TPI spectra for the resultant
29at.% sample are displayed in Fig 3. When these are aligned below 0.8
eV of incident photon energy, we see that in the bandtail region
(between 1.2 and 1.4 eV) the normalized TPI spectra is approximately 7
times the band edge signal of the TPC. As a result, we estimate a hole
collection efficiency of roughtly 75% for this 29at.%Ge sample. This
compares quite well with many PECVD samples studied previously. Although
it is somewhat smaller than the collection fractions found in some of
the best PECVD grown a-Si,Ge:H samples of similar optical gaps (which
can exceed 90% under similar measurement conditions \[9\]), we note that
the sample deposited onto the c-Si substrate exhibited a slightly
broader bandtail than the best HWCVD samples that we had received
previously.

Finally, we used DLCP measurements to examine the defect creation and
annealing kinetics of the highest quality 29at.% HWCVD grown sample and
we compared this to the behavior of a PECVD grown 35at.% Ge sample
(obtained many years ago from United Solar) with a similar optical gap
(near 1.5eV). In both cases we used a tungsten-halogen light source
filtered with a 610 nm long pass filter at a net intensity of roughly
800 mW/cm^2^ (after correcting for absorption in Pd top contact). During
exposure the sample temperatures were maintained below 330K. After each
consecutive period of light soaking the defect densities were determined
by carrying out DLCP at 360 K and 1100 Hz. Figure 4(a) displays one full
set of spatial profiles obtained in this fashion at 1.1 kHz, and Fig.
4(b) displays on log-log axes the spatial averages of the relatively
flat portion of these DLCP profiles obtain the characteristic defect
density as a function of *accumulated* light soaking time. We note that
the defect density rises slowly for the first 3 minutes of light
soaking, reaches a relatively constant plateau, and then subsequently
increases much more rapidly after about 1.5 hours of light soaking.

  --------------------------------------------------- ----------------------------------------------------
  ![](media/image4.wmf){width="3.651388888888889in"   ![](media/image5.wmf){width="2.7493055555555554in"
  height="2.3444444444444446in"}                      height="2.247916666666667in"}

  **FIG. 4(a).** Example of DLCP profiles for the     **FIG. 4(b).** Time evolution of the spatially
  highest quality HWCVD 29at.% Ge alloy sample at     averaged defect density during light soaking
  different stages of light induced degradation.      obtained from Fig. 4(a). These data have been
  Spatial averages of N~DL~ were obtained near 0.55   normalized to an initial State A value of
  μm to construct the time evolution of defect        5.7 × 10^15^ cm^-3^
  density shown in Fig. 4(b).                         
  --------------------------------------------------- ----------------------------------------------------

In Fig. 5(a) we compare the DLCP deduced deep defect densities for the
HWCVD and PECVD a-Si,Ge:H alloys of 1.5eV optical gaps. The densities of
deep defects for both samples have been normalized to their respective
initial State A values. For the PECVD sample, the light induced defect
density also exhibits an initial slow rise and then also shows a much
more rapid increase (beyond roughly the ½ hour mark). The plateau region
is not so distinct for the PECVD sample, while the final rate of
increase of defect density beyond 4.5 hours is actually higher for the
PECVD sample than that of HWCVD sample. The long time regime exhibits
roughly a power law dependence for both samples (with exponents of 0.24
± 0.014 and 0.14 ± 0.012 for the PECVD and HWCVD samples, respectively).

The annealing kinetics for the two samples are compared in Fig. 5(b). In
both cases we employed a series of 10 minute isochronal anneals
beginning at 360K, and increasing in 10K steps up to final anneal at
470K. The same conditions were employed for determining the defect
densities from the DLCP measurements (360K at 1100Hz), and the defect
densities remain normalized to their *initial* State A values.

For the PECVD a‑Si,Ge:H one observes that the defect density decreases
slowly but monotonically, ultimately reaching its initial State A
density after the 470K anneal within experimental error. This is
distinctly different than the behavior of the HWCVD grown sample which
anneals within a much smaller temperature range. This suggests a much
narrower range of annealing energy barriers for this HWCVD sample.
However, at temperatures higher than 430 K the annealing rates appear
nearly same for both the samples.

The data presented above using TPC and DLCP measurements indicate that,
while the low temperature HWCVD grown a‑Si,Ge:H alloys display
electronic properties rivaling those of the best PECVD grown a‑Si,Ge:H
samples, the oxygen impurity level can greatly influence the properties
of the HWCVD alloy samples. That is, it appears that quite low oxygen
levels are required to produce the highest quality HWCVD a-Si,Ge:H
alloys. We also observed a roughly 75% minority carrier collection
efficiency in one 29 at.% Ge alloy sample deposited onto a p^+^ c-Si
substrate. A comparison of the metastable defect creation in HWCVD and
PECVD alloys having similar band gap indicated that a characteristic
two-step degradation process was exhibited in the time dependence of
defect creation for both types of a-Si,Ge:H samples, but this somewhat
more prominent in the HWCVD grown alloys. Subsequent isochronal
annealing experiments showed a distinctly different nature of metastable
recovery of defect densities in HWCVD and PECVD grown a-Si,Ge:H alloys.

  ---------------------------------------- ----------------------------------------------------
  ![](media/image6.wmf){width="3.3125in"   ![](media/image7.wmf){width="3.1729166666666666in"
  height="2.7354166666666666in"}           height="2.6979166666666665in"}

  **FIG. 5(a)**. Comparison of time        **FIG.** **5(b)**. Comparison of the decrease of
  evolution of defect densities during     deep defect densities as a result of a series of 10
  light soaking for the 29at.% Ge HWCVD    minute isochronal anneals at increasing temperatures
  Sample and for a 35at.% Ge United Solar  as indicated. These data suggest that the HWCVD
  PECVD alloy sample. These data have been sample possesess a much narrower range of energy
  normalized to the initial State A        activation barriers for annealing than the PECVD
  values: 5.7 × 10^15^ cm^-3^ for the      sample.
  HWCVD sample and 2.5 × 10^15^ cm^-3^ for 
  the PECVD sample.                        
  ---------------------------------------- ----------------------------------------------------

In addition to the a‑Si,Ge:H alloy studies, we are also continuing our
work on the nc‑Si:H narrow gap materials that are being developed at
United Solar. It now appears that, by incorporating p-i-n nc‑Si:H as the
bottom cell in triple junction, record efficiencies in thin-film Si
based PV will soon be achieved. United Solar (B. Yan) sent us a series
of 6 nc‑Si:H devices early in our new Subcontract period. Three of these
were deposited by MVHF at a rate of 6Å/s, and the other 3 were deposited
by RF at a somewhat lower rate. The MVHF samples were all n‑i‑p devices
deposited onto SS coated with a textured Ag/ZnO back reflector and were
finished with 0.05cm^2^ ITO top contacts. Two of these devices exhibited
relatively high levels of degradation (about 15%) while the third
exhibited very little degradation. The three RF nc‑Si:H samples were
fabricated in three different device structures: One was a n-i-p device
with a Ag/ZnO textured back reflector, one was a n-i-p device on bare
stainless steel, and one was a sandwich device: SS/a‑Si:H/nc‑Si:H/a‑Si:H
where the a‑Si:H capping layers are 0.21μm thick. For all 6 devices the
nc‑Si:H layer was just under 1μm thick. The characteristics of these
samples are summarized in Table I. The goal for this series of samples
was to find out whether we could distinguish different electronic
properties for the high degradation and low degradation films, and
whether we could obtain consistent results on the RF samples independent
of their different device structures.

  --------------------------------------------------- --------------- ----------------------- --------------------------------------------------- ---------------- -------
  **Table I.** United Solar nc‑Si:H sample device                                                                                                                  
  series.                                                                                                                                                          

  **Sample Number**                                   **Substrate**   **Structure**                                                               **Comment**      

  2B 10505                                            SS/Ag/ZnO       n‑i‑p, all nc‑Si:H                                                          Low Degradation  

  2B 10514                                            SS/Ag/ZnO       n‑i‑p, all nc‑Si:H                                                          High Degradation 

  2B 10521                                            SS/Ag/ZnO       n‑i‑p, all nc‑Si:H                                                          High Degradation 

  RF 13512                                            SS/Ag/ZnO       n‑i‑p, all nc‑Si:H                                                          Low Degradation  

  RF 13510                                            Specular SS     n‑i‑p, all nc‑Si:H                                                          Same as 13512    

  RF 13511                                            Specular SS     a‑Si:H/nc‑Si:H/a‑Si:H                                                       Same as 13512    

  ![](media/image8.wmf){width="3.251388888888889in"                                           ![](media/image9.wmf){width="3.251388888888889in"                    
  height="2.6277777777777778in"}                                                              height="2.645138888888889in"}                                        

  **FIG. 6(a)**. DLCP profiles for the high                                                   **FIG.** **6(b)**. DLCP profiles for the same                        
  degradation device 2B10514 in its annealed state.                                           sample in a light-degraded state. We observe almost                  
  The 325K profile contains the largest contribution                                          no difference in the 325K profile, indicating very                   
  from the deep defects, and indicates a quite low                                            little change in the concentration of the deep                       
  value. The rapid rise in the profiles beyond about                                          defects. The lower temperature profiles have been                    
  0.6 microns is probably an artifact caused by the                                           decreased a little, probably due to a slightly                       
  edge of the depletion region encountering the n+                                            deeper Fermi level after prolonged light exposure.                   
  layer at the substrate junction.                                                                                                                                 
  --------------------------------------------------- --------------- ----------------------- --------------------------------------------------- ---------------- -------

Our measurements on this sample series have not been as easy to carry
out compared to the previous set of nc‑Si:H sample devices we obtained
from United Solar. Thus we can only discuss some preliminary findings
for this report. One reason for our difficulties can be understood from
the data exhibited in Fig. 6. In this Figure we compare DLCP profiles
for the high degradation MVHF device 2B10514 in State A and in a
light-degraded State B (obtained by exposure to red-filtered ELH light
at 800mW/cm^2^ for 100 hours). Differences are subtle at best. The
highest temperature profiles, containing the largest contribution from
the deeper defects, are essentially identical. Larger differences
between states A and B are exhibited in the lower temperature profiles,
and this most likely indicates a shift in the Fermi-level position
slightly deeper in the band for the light degraded state.

Figure 6 also indicates that the defect densities are quite low in these
samples compared to those studied previously. Because of this, even
moderate levels of reverse bias (a maximum of 1 volt in this case) tend
to totally deplete these samples as can be seen by the rapid rise in
DLCP densities near 0.8 microns. The higher densities beyond 0.6 μm
cannot be taken at face value; rather, the rapid rise simply reflects
the fact that the edge of the depletion region can no longer move in
response to changes in the device bias. This problem with the sample
thicknesses leads to even more serious issues in interpreting the
transient response measurements, such as the transient photocapacitance
measurements discussed below. For the case of the DLCP measurements,
such as those exhibited in Fig. 6, it means we need to restrict our
discussion to the spatial region smaller than about 0.6 microns.

In Fig. 7 we compare the DLCP curves for the higher degradation sample,
2B10514, with a low degradation sample, 2B10505, both in their annealed
states. It is rather remarkable how similar these curves appear for the
two lower temperature measurements, indicating almost identical carrier
densities for each. At 300K, however, the density for 2B10514 is
significantly higher, and this indicates a higher deep defect density
for the film that exhibits the higher degree of degradation.

  -------------------------- -----------------------------------------------------
  **FIG. 7.** Comparison of  ![](media/image10.wmf){width="3.9381944444444446in"
  DLCP curves for the higher height="2.901388888888889in"}
  degradation sample         
  (2B10514) with a lower     
  degradation sample         
  (2B10505), both in their   
  annealed states. The cell  
  structure was identical    
  and the same range of      
  reverse bias (0 to ‑1V)    
  was employed for both      
  samples. The close match   
  at the lower temperatures  
  implies a nearly equal     
  effective carrier density  
  in both samples, while the 
  deviation at higher        
  temperature indicates a    
  higher deep defect density 
  for the 2B10514 sample     
  device.                    

  -------------------------- -----------------------------------------------------

Finally, in Fig. 8, we give one example of a series of photocapacitance
spectra at different measurement temperatures obtained for sample device
2B10514. As with our reported TPC spectra on the previous series of
nc‑Si:H devices, we observe a decrease in TPC signal above the Si
bandgap as the measurement temperature is increased. This is due to the
increasing fraction of hole collection with increasing temperature.
Indeed, at 325K it appears that only the defect portion of the spectrum
remains visible. We also observed that the hole collection fraction was
significantly better for the annealed state spectra (not shown) compared
to those taken after prolonged light exposure.

However, in contrast to those earlier reported spectra we did not
observe an a‑Si:H component to these spectra at higher measurement
temperatures. Also, the defect portion of the spectrum is nearly 100
times larger than for those observed in the previous series of samples.

We believe that these differences arise from the insufficient thickness
of the i-layers in the current series of samples. That is, during the
capacitance transient the edge of the depletion region expands beyond
its position under ambient reverse bias. Because these i-layers are
already nearly completely depleted under small values of reverse bias,
this means that the edge of the depletion region extends into the n^+^
layer near the back contact after the filling pulse and during the
transient. Moreover, the region near the edge of the depletion region
contributes more strongly to the TPC signal than the interior region.
This means that the spectra shown in Fig. 8 are probably dominated by
the n^+^ nc‑Si:H layer near the substrate contact. This would account
both for the larger defect magnitude and the absence of visible a‑Si:H
component.

This situation also makes it nearly impossible to learn anything from
the TPC spectra regarding the difference in electronic properties
between the high degradation and low degradation nc‑Si:H material from
the current series of samples. Fortunately a new series of nc‑Si:H
sample devices is now scheduled for preparation at United Solar which
are to be fabricated with somewhat thicker i-layers.

  -------------------------- -----------------------------------------------------
  **FIG. 8.**                ![](media/image11.wmf){width="3.3048611111111112in"
  Photocapacitance spectra   height="3.8152777777777778in"}
  for sample device 2B10514  
  in its light degraded      
  state. Spectra were taken  
  at 1.1kHz with the ambient 
  reverse bias at --0.2volts 
  and filling pulse height   
  of 0.3volts (to            
  +0.1volts). The signal     
  magnitude decreases above  
  the Si bandgap with        
  increasing temperature due 
  to the increasing fraction 
  of holes collected. The    
  defect band signal is      
  nearly temperature         
  independent and much       
  larger in magnitude than   
  for the previous series of 
  United Solar nc‑Si:H       
  samples studied. This is   
  probably due to an         
  insuf­ficient thickness of  
  the i-layer in these       
  sample devices, so that    
  the spectra are dominated  
  by the properties of the   
  material near the back     
  contact; that is, in the   
  n^+^ region of the sample  
  device.                    

  -------------------------- -----------------------------------------------------

As we have discussed previously, the variation in the appearance of the
photo-capacitance spectra indicates the mixed phase nature of the United
Solar nc-Si:H material. Specifically, we believe that the progression of
these spectra from being more μc-Si:H like at lower temperatures to more
a‑Si:H like at higher temperatures is due to the suppression of the
crystallite signal component when *more of the minority carriers are
able to escape* the depletion region during the TPC measurement time
window. This cancels the charge change (and hence the photocapacitance
signal) caused by the escape of the optically excited majority carriers
that dominates the appearance of the spectra at the lower temperatures.
In contrast, the minority carriers generated in the amorphous silicon
component of the sample appear to remain trapped and so this component
of the sample dominates the spectra above 280K.

This means, in particular, that the *higher* the magnitude of the TPC
signal near 1.5eV (where the microcrystalline Si absorption tends to
dominate) the lower the hole collection fraction will be for carriers
generated in the microcrystalline component of this nc-Si:H sample. In
our previous Quarterly, we reported that prolonged light exposure also
reduces the fraction of holes collected. In Fig. 2 we display two sets
of spectra for sample 14657. In Fig. 2(a) we show the effect of reducing
the temperature from 240K to 180K. We observe that the TPC signal near
1.5eV increases by more than a factor of 30, indicating that the hole
collection fraction has been substantially reduced at 180K compared to
240K. In Fig. 2(b) we display the effect on the 240K TPC spectrum of
light soaking for 100 hours to 610nm filtered ELH light at an intensity
of 100mW/cm^2^. The resulting spectrum is quite similar to that obtained
by lowering the measurement temperature, indicating a similar decrease
in hole collection.

In spite of the clear reduction in hole collection, Fig. 2(b) indicates
that the light exposure does *not* appear to increase the deep defect
density (dangling bonds). Thus, the root cause of the decrease in hole
collection is not yet understood.

To try to gain some additional insight into the cause of the reduced
hole collection, we have determined the magnitude of the TPC signal as a
function of temperature near 1.5eV in more detail. This is plotted in
Fig. 3 in Arrhenius form for both the annealed (State A) and light
degraded state (State B) of this sample. As we see, for both cases the
hole collection fraction vs. temperature appears to be thermally
activated, with characteristic energies near 0.2eV in both cases. This
may therefore suggest a predominant hole trap near the valence band edge
which increases in magnitude after prolonged light exposure. This trap
does not appear to be simply related to the dangling bond defects in the
a‑Si:H component of the sample, although it could still originate from
the a‑Si:H component of these mixed phase samples, or from the boundary
region between the two phases. This suggests the potential value of
future experiments which compare the degree of light induced effects for
nc‑Si:H samples which contain varying fractions of the a‑Si:H component.

![](media/image10.wmf){width="4.239583333333333in"
height="3.5520833333333335in"}![](media/image11.wmf){width="3.3020833333333335in"
height="4.020833333333333in"}

[^1]: . C. E. Michelson and J.D. Cohen, Phys. Rev. B**41** (1990) 1529.