**The Effect of Mo Deposition Conditions on Defect Formation and Device
Performance for CIGS Solar Cells**

V. Mohanakrishnaswamy, H. Sankaranarayanan, S. Pethe, C. S. Ferekides
and D. L. Morel

Department of Electrical Engineering

Clean Energy Research Center

University of South Florida

Tampa, FL

**ABSTRACT**

We have studied the effect of native defects on CIGS solar cells on an
ongoing basis and recently investigated the role of Na. We modulated
access of Na to the growing film by use of Si~3~N~4~ blocking layers as
well as by varying the Mo thickness and density. Our process is tuned to
the level of Na coming from our soda lime glass substrates, however, the
role of Na is complex. Collectively the results suggest that one role of
Na is that of a catalytic agent for the oxidation of V~Se~, which is a
donor-like defect, to acceptor-like O~Se~. In simulating the effect of
this defect in AMPS we find that its biggest effect on performance is
its presence in the junction interface region, and its presence there
gives rise to the Jsc -- Voc trade-off that we observe in our
experimental data. This results in the efficiency being pinned in the 13
-- 14% range for our deposition process and explains the difficulties we
have had in overcoming this mechanism. These insights also indicate the
path to overcoming this performance limitation.

**BACKGROUND**

We have been working to improve understanding of the role of defects in
CIGS devices. Our approach is based upon use of defect properties
generated by the NREL theory group as input for AMPS© device simulation.
Comparisons are made between simulations and experimental data to verify
self-consistency. We have determined that a combination of V~Cu~ defects
and (M~Cu~ + 2V~Cu~ ) defect pairs can explain much of our experimental
data. However, there are inconsistencies that need to be addressed.
While these two defects play an important role, it is clear that other
defects are operative. Recently we reported\[1\] on the effects of
various impurities in our devices. We primarily focused on impurities
attributable to the process chamber environment, but there were also
preliminary results on contributions from the substrate, particularly
Na. During subsequent experimentation we learned that substrate effects
are much stronger than thought, and we present our findings in this
paper. As will be shown, the results provide significant insights to the
effect of Mo substrate properties on device performance and clearly
indicate what is limiting performance.

**EXPERIMENTAL**

Our CIGS layers are deposited by the all-solid-state two-step process.
While overall Ga content is in the 10 -- 20% range, only a fraction
enters the space charge region. Consequently, device band gaps are
typically 0.95 -- 1.0 eV. We use two processing systems to determine the
influence of the deposition environment on performance. The first is a
standard bell jar chamber without a load lock. The second has a load
lock and two deposition chambers to allow isolation of Se background
flux during metal deposition. The standard deposition procedure in this
system is as follows. Cu and Ga are deposited sequentially in the first
chamber. The substrate is then transported to the second chamber(under
vacuum) in which In is deposited in the presence of Se flux.
Selenization is completed by leaving the Se flux on while ramping up to
the final deposition temperature of 550 ºC. During this ramp a 25 Å
layer of Cu is added. Overall the substrate is exposed to Se flux of
about three times that needed to form a stoichiometric film. The CIGS
deposition process takes about 30 minutes. The process described for the
two-chamber system is the same as that used in the single chamber. The
important difference is that the Cu and Ga precursor layers are
deposited in a Se-free environment in the two-chamber system.

Junction formation is by the standard CBD CdS and sputtered ZnO
techniques. A key process step for these studies is the Mo layer. It is
deposited by DC magnetron sputtering using the standard high and low
pressure techniques to meet adhesion and conductivity requirements. As
will be discussed, Mo properties are sensitive to the details of the
deposition parameters. We have intentionally varied these parameters as
a way of altering the Mo properties, particularly with respect to its
influence on Na and other impurities entering the growing film from the
glass substrate.

For one part of the study layers of Si~3~N~4~ are deposited onto the
glass substrate by RF magnetron sputtering from a Si~3~N~4~ target prior
to Mo deposition. The Si~3~N~4~ layer thickness is varied to control the
flow of impurities from the sodalime glass substrate. Additional
relevant details of the experimental procedures will be presented below.

**RESULTS AND DISCUSSION**

The fact that Na plays a significant role in device formation and
performance is well-known. We are particularly interested in determining
its influence on defect formation and the effect of those defects on
performance. To this end we conducted a series of experiments to
modulate its entry to the CIGS layer. While the techniques used do have
possible additional influence on film growth beyond Na-induced defect
formation, the discussion here will focus predominantly on that aspect.
Also, preliminary results on structural changes in the CIGS films due to
the substrate effects indicate that in most cases they are secondary to
the influence of Na on point defects. Nevertheless, efforts to
understand these effects are continuing as they hold further sway over
device performance.

From the beginning of this project we have tried to limit ourselves to
processing procedures that could be practiced in industry. Along these
lines we have used ordinary window glass as our substrate.
Unfortunately, all such glass is not the same, and we have found our
process to be sensitive to starting new glass lots. Some of the
difficulties are attributable to scratches in the glass, and we have
worked around those and consider them to be manageable. However, there
are more subtle effects that have their origins in the composition,
density and other properties of the glass. In undertaking these studies
we sought two things: 1. to understand the mechanisms influenced by the
substrate, and 2. to perhaps develop an acceptable pre-coating procedure
that would render all glass equivalent.

![](media/image1.wmf){width="3.25in"
height="2.511111111111111in"}~3~N~4~ is commonly used as a blocking
layer, so we prepared a series of devices with various thicknesses of
Si~3~N~4~ deposited on the glass prior to Mo deposition. In Fig. 1 we
show the effect of a Si~3~N~4~ blocking layer on Voc. With increasing
Si~3~N~4~ thickness less Na enters the CIGS film causing Voc to
diminish. Jsc behaves similarly(not shown). In Fig. 2 is shown the
effect of the Si~3~N~4~ layer on doping concentration as determined from
the reverse bias slope for CV measurements. The similarity of these
plots suggests that the net acceptor concentration is affected by Na
content, and that correlates with Voc behavior. Since the V~Cu~ defect
has the lowest formation energy of the acceptor-like defects\[2\] and it
is shallow, we assign it as the predominant acceptor. One role of Na
might be as a substitutional dopant for Cu. If so, the V~Cu~ defects
would diminish with increasing Na which is opposite to what is observed
in Fig 2. Also we find with AMPS simulations that just increasing V~Cu~
acceptors to simulate increasing Si~3~N~4~ thickness does not produce
Voc (and Jsc) behavior consistent with the data of Fig. 1. To achieve
that consistency we find it necessary to include a compensating donor
that is affected by Na. Our choice in this case is the V~Se~ defect
which is donor-like at Ec -- 0.08 eV\[2\]. And we adopt the model that
has Na in the role of a catalyst that promotes oxidation of the V~Se~
that converts it to acceptor-like\[3\]. Thus the net doping
concentration is due to compensation between V~Cu~ and V~Se~ defects. As
Na is reduced by the blocking layer, less oxidation occurs and more
V~Se~ donors remain to compensate the V~Cu~ acceptors.

Another aspect of the data in Figs 1. and 2. is the more rapid fall of
Voc for the single chamber. We attribute this to the different
processing environment of the chambers. In particular, the presence of
Se background flux in the single chamber environment gives rise to
formation of Cu~x~Se~y~. This is known to increase the conductivity and
hurt junction performance, although Voc's in the absence of a Si~3~N~4~
blocking layer are almost comparable to those from the two-chamber
environment. This species is formed during Cu deposition and is
difficult to eliminate once it is formed. The more rapid fall of Voc in
the single chamber environment suggests that the oxidation mechanism
discussed above is influenced in some way by the presence of this
species. It would seem to play a role in preventing the oxidation of
V~Se~ defects and thus unfavorably altering the balance of compensation.
Further experimentation will be necessary to sort this out

In the next series of experiments the Si~3~N~4~ layer was eliminated and
the thickness of the Mo layer was varied. These devices were made in the
single chamber system. By making the Mo layer thinner than the standard
thickness of 10,500 Å we could allow more Na to enter the CIGS growth
area. The results are shown in Fig. 3. As can be seen, as the Mo
thickness is decreased, Jsc rises slowly. This is essentially a
continuation of the trend from the Si~3~N~4~(not shown) series in which
Jsc increased with decreasing Si~3~N~4~ thickness. While there is a
real, though subtle effect of Na on Jsc, the effect on Voc is more
dramatic. In Fig. 1 we see a rapid rise in Voc with decreasing
Si~3~N~4~![](media/image3.wmf){width="3.375in"
height="2.933333333333333in"} thickness that we have attributed to Na.
In Fig. 3 we see Voc starting to drop slowly at first as we decrease the
Mo thickness, and then the drop accelerates in the 4000--5000 Ǻ range.
Taken together these results indicate that our process is tuned to the
amount of Na provided by a 10,500 Å thick Mo layer. Any deviation from
this Na level results in drop off in Voc. However, the slow rise in Jsc
suggests a trade-off that we must optimize. This will be discussed
further below.

![](media/image4.wmf){width="3.25in" height="2.796527777777778in"} Above
we have suggested that Na does not appear to play a substitutional role
for Cu. However, a closer examination of our data suggests that it does
influence achievement of an optimized metal ratio. This can be seen in
Fig. 4 which shows the Voc dependence versus metal composition for two
different Mo layer thicknesses. Because we create an intentional metal
gradient on our substrate during deposition we are able to see
compositional effects on performance with data from a single deposition.
The x axis in Fig. 4 represents a row of 5 devices in which device 1 has
the highest Cu/In ratio, and device 5 the lowest. The ratio range is
about 0.9 -- 1.0. As can be seen, Voc dependence on this ratio for a Mo
thickness of 6000 Å is much less than that for 10,500 Å. This is
important from a commercialization perspective to improve yield.
However, we must be mindful of the data in Fig. 3 that indicates that at
6000 Å Mo thickness we are at the edge of a big drop in Voc. If our Mo
thickness strays a bit to the thin side, the magnitude of Voc could drop
sharply. This is yet another example of the many trade-offs that must be
managed in such a complex deposition space.

![](media/image5.wmf){width="3.25in" height="2.0909722222222222in"}^-4^
Ω-cm at a rate of 1.3 Å/s to 8.0 x 10^-5^ Ω-cm at 3.3 Ǻ/s as shown in
Fig. 5. The effect of the rate on Voc and Jsc is shown in Fig. 6. We
expect some contributions from Na effects since lower resistivity is in
part due to higher Mo density, and thus as the rate increases and
densifies the Mo layer, we would expect lower Na penetration. However,
there are other factors at work here as well. The data was taken with a
newer Mo target, and the higher rates were accompanied by higher target
voltage(\> 600 volts at 3.3 Å/s). The open data points in figure 4 were
taken at the same high rate of 3.3 Å/s, but with the target broken in,
and thus required only 400 volts for the target voltage. Thus both Jsc
and Voc are affected by target voltage in addition to effects
attributable to Mo densification. It is also apparent that many of these
effects result in Voc and Jsc moving in opposing directions.

We are pursuing further understanding of both the Na-driven as well as
structural influence of substrate properties on film growth. We have
measured the Na content in our films using SPS and find qualitative
correlation with expectations based upon the Na controlling mechanisms
used. However, we still do not have reliable enough quantitative data to
present at this point. We will continue to pursue this matter as well as
the other experimental issues raised by the data presented above.

![](media/image6.wmf){width="3.3125in" height="2.75in"}

While continuing to add to our experimental data set we have also been
attempting simulations of our suggested mechanisms in AMPS. Using the
model of Na catalyzed oxidation of V~Se~ donor defects in AMPS we
observe good consistency with experimental results. To reiterate the
major aspects of our proposed model: V~Cu~ 's are the primary acceptor
states that are shallow enough to be completely ionized under most
conditions. While effects due to (M~Cu~ + 2V~Cu~ ) defect pairs can also
be operative\[1\], we propose that V~Se~ donor defects play a
significant role with respect to Na. These are oxidized to acceptor-like
states in the presence of Na. The net compensation among these
determines the location of the Fermi energy and performance. These
compensation effects seems to be most influential in the top surface
region of devices. In Fig. 7 we show the AMPS simulated effect on device
performance of the presence of donor defects in the surface region. The
defect is located at Ec -- 0.08 eV to simulate V~Se~'s\[2\]. As can be
seen, lowering of the defect density helps Voc, but at the expense of
Jsc, in good agreement with the experimental results presented above.
The net effect at the performance level, however, is "pinning" of the
efficiency at the 14% level. This is commonly observed in our data, and
now we understand that such observations and limitations to performance
may be attributable to substrate properties and the influence of those
properties on Na migration. This is a major new insight to the
performance of our devices and should provide new pathways for
overcoming these limitations.

**CONCLUSIONS**

We have gained considerable insights to the influence of substrate
properties on device performance. A key factor is the control of Na to
the CIGS film growth area. A given deposition process must be tuned to
the details of Na availability. The availability in a given process can
be influenced by parameters such as Mo thickness and density as well as
sputtering voltage which is influenced by break in of the target. While
Na can influence film properties in a structural way, we have focused on
the defect nature of its influence. We have proposed that V~Se~ donor
defects play a significant compensating role and thereby influence
device performance. Further we have suggested that Na helps catalyze the
oxidation of these states. In so doing they become acceptors which
results in a shift in compensation balance. This effect seems to be
strongest in the top surface region of our devices. As a result of the
dynamics of compensation Voc and Jsc show opposing tendencies which
results in pinning of efficiency, for our devices, at the 14% level. We
have observed this effect for some time and now realize its origins. The
challenge now is to use this understanding to decouple Voc and Jsc and
move the efficiency forward.

*Acknowledgements:* This project has been funded by the NREL/DOE
Thin-Film Partnership.

*AMPS* is a first principles simulation code developed by Penn
State/EPRI.

**REFERENCES**

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