A national laboratory of the U.S. Department of Energy

Office of Energy Efficiency & Renewable Energy

National Renewable Energy Laboratory 

Innovation for Our Energy Future 

Fundamental Research and 

Development for Improved 

Crystalline Silicon Solar Cells 

Final Subcontract Report 

March 2002 — July 2006 

A. Rohatgi 

Georgia Institute of Technology 

Atlanta, Georgia 

Subcontract Report 

NREL/SR-520-42324 

November 2007 

NREL is operated by Midwest Research Institute ● Battelle Contract No. DE-AC36-99-GO10337 
National Renewable Energy Laboratory

1617 Cole Boulevard, Golden, Colorado 80401-3393 

303-275-3000 • www.nrel.gov 

Operated for the U.S. Department of Energy 

Office of Energy Efficiency and Renewable Energy 

by Midwest Research Institute • Battelle 

Contract No. DE-AC36-99-GO10337

Subcontract Report 

NREL/SR-520-42324 

November 2007 

Fundamental Research and 

Development for Improved 

Crystalline Silicon Solar Cells 

Final Subcontract Report 

March 2002 — July 2006 

A. Rohatgi 

Georgia Institute of Technology 

Atlanta, Georgia 

NREL Technical Monitor: R. Matson/F. Posey-Eddy 

Prepared under Subcontract No. AAT-2-31605-02 
This publication was reproduced from the best available copy 

submitted by the subcontractor and received no editorial review at NREL

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Table of Contents

Summary ....................................................................................................................................... v

1. PECVD SiN-Induced Defect Passivation and Rapid Thermal Processing for HighEfficiency EFG Silicon Solar Cells......................................................................................... 1

1.1 Hydrogen Content and Defect Passivation Capability of Low- and High-Frequency 

PECVD SiN Films.................................................................................................................. 1 

1.1.1 Impact of Surface Pretreatment with NH3 Prior to PECVD SiN Deposition .......... 2 

1.1.2 Impact of As-Grown Bulk Lifetime on SiN-(2)-Induced Defect Passivation of 

Screen-Printed EFG Silicon Solar Cells ............................................................... 4 

1.2 Process Design, Fabrication, and Analysis of High-Efficiency Screen-Printed EFG Si 

Solar Cells............................................................................................................................... 5 

1.2.1 Development of Optimal Process Sequence for EFG Si Solar Cells with HighFrequency PECVD SiN ........................................................................................ 6 

1.2.2 Development of Optimal Process Sequence for EFG Silicon Solar Cells with 

Low-Frequency PECVD SiN................................................................................ 9 

1.2.3 Optimization of Hydrogenation Temperature in RTP for Low-Frequency 

PECVD SiN to Achieve High-Efficiency EFG Si Cells with 

Photolithography Contacts.................................................................................. 10 

1.2.4 Optimization of Hydrogenation Temperature for SiN-(2) in RTP......................... 10 

1.3 Fabrication and Analysis of 15.8% Efficient EFG Si Solar Cell with Photolithography 

Contacts................................................................................................................................. 12 

1.4 Fabrication and Analysis of Record-High Efficiency Screen-Printed EFG Silicon Solar 

Cell........................................................................................................................................ 15 

1.5 Conclusion ............................................................................................................................ 21

2. RTP-enhanced Hydrogen Passivation in Ribbon Multicrystalline Silicon Studied by 

Scanning Room-Temperature Photoluminescence Spectroscopy......................................... 22 

2.1 Introduction........................................................................................................................... 22 

2.2 Solar cell fabrication and photoluminescence analysis procedures...................................... 22 

2.3 Characterization of defects in EFG Si before and after hydrogenation by room 

temperature PL mapping....................................................................................................... 24 

2.4 Conclusions........................................................................................................................... 27 

2.5 References............................................................................................................................. 28

3. Effect of Material Inhomogeneity on the Open-Circuit Voltage of String Ribbon Si Solar 

Cells ...................................................................................................................................... 29 

3.1 Introduction........................................................................................................................... 29 

3.2 Experiment............................................................................................................................ 31 

3.3 Results and discussion .......................................................................................................... 31 

3.3.1 LBIC scans and IQE measurements....................................................................... 31 

3.4 Extraction of effective diffusion length Leff from the IQE response..................................... 33 

3.5 Theoretical and experimental assessment of the impact of electrically active defects on 

solar cell performance ........................................................................................................... 36 

3.5.1 Development of the analytical model to assess the loss in VOC resulting from 

inhomogeneity..................................................................................................... 36 

iii
3.5.2 Model calculations to assess the loss in VOC of a cell with two regions of 

different recombination intensity........................................................................ 38 

3.5.3 Model calculations to assess the loss in VOC of a cell with three regions of 

different recombination intensity........................................................................ 39 

3.5.4 Application of the analytical model to defective cells........................................... 40 

3.6 Conclusions........................................................................................................................... 42 

3.7 References............................................................................................................................. 43

4. Resistivity and Lifetime Variation along Commercially Grown Ga- and B-Doped 

Czochralski Si Ingots and Its Effect on Light-Induced Degradation and Performance of 

Solar Cells............................................................................................................................. 45 

4.1 Introduction........................................................................................................................... 45 

4.2 Experimental procedures....................................................................................................... 46 

4.3 Results and discussion .......................................................................................................... 46 

4.3.1 Growth of Ga-doped Cz ingots.............................................................................. 46 

4.3.2 Resistivity distribution along commercially grown B-doped and Ga-doped Cz 

ingots................................................................................................................... 46 

4.3.3 As-grown and post-diffusion lifetime in wafers from B- and Ga-doped Cz 

ingots................................................................................................................... 47 

4.4 LID and performance of screen-printed solar cells............................................................... 49 

4.4.1 B-doped ingots....................................................................................................... 49 

4.4.2 Ga-doped ingot....................................................................................................... 53 

4.5 Conclusions........................................................................................................................... 54 

4.6 References............................................................................................................................. 55

5. Understanding the Formation of High-Quality Thick-Film Ag Contacts on High SheetResistance Si Emitters for Solar Cells .................................................................................. 57 

5.1 Introduction........................................................................................................................... 57 

5.2 Experimental method ............................................................................................................ 57 

5.3 Results and discussion .......................................................................................................... 58 

5.3.1 Effect of Firing Temperature on the Ag-Si Contact Interface for PV168 Paste .... 59 

5.3.2 Investigation of Screen-Printed Contacts to 100 Ω/sq Emitter Using Widely 

Used Commercial Pastes A and B and Conventional Firing Temperature......... 65 

5.3.3 High Temperature (835°C) Firing of Conventional Pastes A and B on the 100 

Ω/sq Emitter........................................................................................................ 67 

5.4 Conclusion ............................................................................................................................ 69 

5.5 References............................................................................................................................. 70 

 

iv
Summary 

This report summarizes the progress made by Georgia Tech in the 2002-2006 period 

towards high-efficiency low-cost crystalline silicon solar cells under NREL Project 

# RCQ-1-31605 entitled Development of Low-Cost High-Efficiency Silicon Cells Through 

Fundamental Understanding of Critical Material and Technology Issues”. The overall goal 

of the program is to improve the cost effectiveness of crystalline silicon photovoltaics by 

advancing the current understanding of hydrogen passivation and removal of harmful 

defects in silicon along with the development of novel conductive pastes and technologies 

for low resistance screen printed contacts. This program emphasizes fundamental and 

applied research on commercial substrates and manufacturable technologies. A 

combination of material characterization, device modeling, technology development, and 

complete cell fabrication are used to accomplish the goals of this program. This report is 

divided into five sections that summarize our work on i) PECVD SIN-induced defect 

passivation (Sections 1 and 2); ii) the effect of material inhomogeneity on the performance of mc-Si 

solar cells (Section 3); iii) a comparison of light induced degradation in commercially grown Gaand B-doped Czochralski Si ingots (Section 4);and iv) the understanding of the formation of highquality thick-film Ag contacts on high sheet resistance emitters (Section 5). 

In Section 1, we investigate bulk defect hydrogenation via PECVD SiN deposition and 

anneal. We have found that a PECVD SiN film deposited at a frequency of ~150 KHz at 

~400 °C is superior to the SiN deposited with a frequency of 13.56 MHz at 300 °C, despite 

the fact that the high-frequency SiN film contains more atomic hydrogen in the form of SiH and N-H bonds. This suggests an additional source of hydrogen for the low-frequency 

SiN film, which is attributed to NH3 pre-treatment during the low-frequency SiN 

deposition. The low-frequency SiN film showed an optimum hydrogenation temperature of 

~750 °C while the high-frequency SiN film showed the optimum at ~800 °C in RTP. A ~2 

μs as-grown lifetime increased to ~25 μs after 850 °C firing and ~50 μs after ~750 °C 

firing of the high- and low-frequency SiN films, respectively. A complete process 

sequence was developed through fundamental understanding to achieve record-high 

efficiency EFG Si cells. RTP contact firing with fast ramp-up and cooling gave ~1.4% increase 

in absolute cell efficiency compared to the slow belt firing used in production. This 

improvement was largely due to bulk lifetime enhancement associated with rapid cooling 

in RTP, which enhances the retention of hydrogen at the defects in EFG Si cells. It was 

found that the RTP contact firing (~700 °C/1 s) of the low-frequency SiN film with fast 

cooling rate (≥ 40 °C/s) could preserve the hydrogenated bulk lifetime due to higher 

retention of hydrogen a the defects. A two-step RTP processes, involving defect 

passivation at ~750 °C/60 s and screen-printed contact firing at ~700 °C/1 s, produced a 

record-high cell efficiency of 15.9% on Si ribbon. 

In Section 2, we analyze the electrical activity of defects in ribbon mc-Si by scanning 

room-temperature photoluminescence spectroscopy before and after hydrogen passivation. 

This analysis revealed a strong inhomogeneity in the increase of minority carrier lifetime 

caused by the hydrogen defect passivation in mc-Si. We present experimental evidence 

that RTP-Al/SiNx processing leads to strong lifetime enhancement caused by hydrogen 

defect passivation in low lifetime regions of mc-Si wafers. Additional details on the 

hydrogenation mechanism are revealed in a course of the de-hydrogenation study. 

Hydrogen out-diffusion shows a different rate or activation energy between high and low 

lifetime regions on the wafers. 

The effect of material inhomogeneity on open-circuit voltage of String Ribbon Si solar 

cells is investigated by a combination of experimental results and a simple analytical 

v
model in Section 3. Light beam-induced current measurements showed that a cell with no 

detectable defective region gave an efficiency of 15.9% with VOC of 616 mV. However, a 

neighboring cell with highly defective regions covering 38% of its area, as determined by 

LBIC measurements, gave an efficiency of 14.1% with VOC of 578 mV. Another cell with 

19% highly defective regions gave an efficiency of 15.0% with VOC of 592 mV. A simple 

and approximate analytical model was developed to quantify the loss in VOC on the basis 

of recombination intensity and the area fraction of defective regions. This model showed 

that the majority of the loss in VOC is associated with the most defective region, even if its 

area fraction is relatively small. 

A systematic study of the variation in resistivity and lifetime on cell performance, 

before and after light-induced degradation in commercially grown B- and Ga-doped Cz 

ingots is presented in Section 4. Manufacturable screen-printed solar cells were fabricated 

and analyzed from different locations on the ingots. Despite the large variation in 

resistivity (0.57 to 2.5 Ω-cm) and lifetime (100-1000 μs) in the Ga-doped Cz ingot, the 

efficiency variation was found to be ≤ 0.5% with an average efficiency of ~17.1%. No LID 

was observed in these cells. In contrast to the Ga-doped ingot, the B-doped ingot showed a 

relatively tight resistivity range (0.87 Ω-cm to 1.22 Ω-cm), resulting in smaller spread in 

lifetime (60-400 μs) and efficiency (16.5-16.7%) along the ingot. However, the LID 

reduced the efficiency of these B-doped cells by about 1.1% absolute. Additionally, the use 

of thinner substrate and higher resistivity (4.3 Ω-cm) B-doped Cz was found to reduce the 

LID significantly, resulting in an efficiency reduction of 0.5-0.6% as opposed to >1.0% in 

~1 Ω-cm ~17% efficient screen-printed cells. As a result, Ga-doped Cz cells gave 1.5% 

and 0.7% higher stabilized efficiency relative to 1 Ω-cm and 4.3 Ω-cm B-doped Cz Si 

cells, respectively. 

Finally in Section 5, the physical and electrical properties of screen-printed Ag thickfilm contacts are studied in an effort to understand and achieve good ohmic contacts to 

high sheet-resistance emitters. A combination of cross-sectional SEM/TEM, AFM, and 

SIMS measurements is used to analyze the Ag-Si interface and explain the electrical 

performance of three different screen-printed Ag pastes (A, B, and PV168) subjected to 

high and low-temperature firing conditions. Contact-resistance measurements and the solar 

cell parameters are used to establish a correlation between the electrical and physical 

properties of screen-printed contacts. Interface analyses showed that Ag does not make a 

full area contact to Si. Instead, there are dispersed Ag crystallites that are in direct contact 

to a small fraction of Si underneath the Ag grid. It was also found that larger Ag 

crystallites are formed at higher firing temperature, thus increasing the Ag-Si contact area 

fraction. At 750°C firing, all three pastes failed on a 100 Ω/sq emitter; PV168 and paste A 

failed because of a much smaller Ag crystallites and Ag-Si contact area fraction at the 

interface, and paste B failed because of excessive Ag penetration into the p-n junction. At a 

higher firing temperature of ~835°C, conventional paste A failed because of excessive 

diffusion of Ag and impurities into the emitter region which degraded the contact quality. 

Paste B failed because of the formation of very large (0.6-1 μm) Ag crystallites, which 

shunted the p-n junction. Of the three pastes, only the PV168 paste from DuPont gave 

acceptable contact quality on a 100 Ω/sq emitter with a solar cell fill factor (FF) of 0.782. 

This was the result of a more uniform distribution of re-grown Ag crystallites at the Ag-Si 

contact interface, in conjunction with significantly reduced Ag penetration into the Si 

emitter region. 

vi
1. PECVD SiN-Induced Defect Passivation and Rapid Thermal 

Processing for High-Efficiency EFG Silicon Solar Cells 

1.1 Hydrogen Content and Defect Passivation Capability of Low- and HighFrequency PECVD SiN Films 

Hydrogenation capability of a PECVD SiN can be a function of the deposition 

conditions such as plasma excitation frequency, deposition temperature, a ratio of NH3 to 

SiH4, or a pre-deposition NH3 treatment. In this report, we investigated two different 

PECVD SiN films for defect passivation of EFG Si solar cells. A PECVD SiN film 

deposited with a plasma excitation frequency of 13.56 MHz at 300 °C is referred to as SiN-

(1) while the SiN-(2) film was deposited at ASE Americas, using a frequency of 150 KHz 

and at temperatures of ≥ 400 °C in a tubular reactor. 

Table 1-1 shows the difference in the important deposition variables for the two 

different SiN films used in this investigation. FTIR measurements were performed on the 

two SiN films, which showed that the SiN-(1) contains ~2.5 times more atomic hydrogen 

than the SiN-(2) after deposition but slightly less after the annealing at 850 °C for 2 min. It 

is important to note that FTIR only detects the atomic hydrogen in the form of Si-H and NH bonds. Since the data in Table 1-1 indicates that much more atomic hydrogen is released 

from the SiN-(1) after 850 °C anneal, one would expect more hydrogenation from the SiN-

(1). 

Table 1-1. Deposition variables and characteristics of two different PECVD SiN films. 

Variable SiN-(1) SiN-(2) 

Excitation frequency 13.56 MHz 150 KHz 

Deposition temperature (°C) 300 ~400 

Thickness (Å) as-deposited ~830 ~780 

Thickness (Å) after anneal ~750 ~750 

H content (cm-3) as-deposited 2.5×1022 1.1×1022

H content (cm-3) after anneal ~5×1021 ~7×1021

Surface pre-treatment None 

NH3 gas flow 

prior to SiN 

deposition 

However, Figure 1-1 shows that this is not the case for the two different heat treatments 

in RTP: 700 °C/1s and 850 °C/120 s. These treatments were established for front contact 

and BSF formation, respectively. It was found that the SiN-(1) produced greater lifetime 

enhancement at 850 °C than 700 °C. On the contrary, the SiN-(2) showed that 

hydrogenation at 700 °C is much more effective than at 850 °C. In addition, the final 

lifetime values for the SiN-(2) passivation are much higher than the SiN-(1) for both the 

annealing conditions. This seems contrary to the FTIR data, which showed greater release 

of hydrogen from the SiN-(1). Furthermore, dehydrogenation seems more important for the 

SiN-(2) than the supply of hydrogen because 700 °C anneal gives superior passivation than 

1
850 °C. This could happen if there is another large source of atomic hydrogen, in addition 

to the Si-H and N-H bonds inside the film, which makes the defect passivation 

dehydrogenation-limited rather than supply-limited. At low frequencies, the ions in the 

plasma can follow the excitation frequency and hence are expected to have higher energy 

and cause more damage to the wafer surface. In addition, the surface pre-treatment with 

NH3 gas during the SiN-(2) deposition can implant hydrogen into the Si surface. 

According to Sopori et al. [NCPV Review Meeting, Colorado, 2001], surface damage in Si 

can absorb (or trap) a lot of atomic hydrogen. This form of hydrogen may not show up in 

the FTIR measurements but may be readily available for release and diffusion. Thus defect 

passivation using the SiN-(2) may not be limited by the supply of atomic hydrogen from 

the SiN film but by the retention of hydrogen at the defects. In order to support this 

hypothesis, a study was conducted in the next section, with and without the NH3 surface 

pre-treatment. 

1.5 1.6

17.0

6.6

38.4

13.8

0

5

10

15

20

25

30

35

40

45

SiN-(1) SiN-(2)

Average Bulk Lifetime ( μs)

As-Grown Lifetime

after 700C/1s anneal

after 850C/2min anneal

Figure 1-1. Comparison of bulk lifetimes in the EFG Si samples coated with two 

different PECVD SiN films. After SiN deposition on the front and Al screenprinting on the back, the samples were annealed simultaneously at 700 °C and 

850 °C in RTP. 

1.1.1 Impact of Surface Pretreatment with NH3 Prior to PECVD SiN Deposition 

A systematic experiment was designed and conducted to understand the role of NH3

surface pre-treatment on the degree of defect passivation of EFG Si. In this experiment, 

EFG Si wafers with the as-grown bulk lifetime of ~2 μs were used. After P diffusion at 

930 °C for 6 min in the belt furnace, the bulk lifetime of the samples increased to ~6 μs 

due to P gettering. Then, the SiN-(2) film was deposited without (Scheme A) and with 

(Scheme B) the surface pre-treatment with NH3 followed by 700 °C anneal in RTP to 

induce hydrogen passivation of defects. 

It was found that annealing of SiN-(2) without the NH3 pre-treatment (Scheme A) 

increased the bulk lifetime to ~18 μs while annealing of SiN-(2) with the NH3 pre2
treatment (Scheme B) raised the bulk lifetime to ~36 μs (Figure 1-2). This clearly shows 

that the NH3 pre-treatment helps the hydrogenation of defects in EFG Si. More research 

needs to show that this is due to hydrogen implantation or absorption of more atomic 

hydrogen at the Si and SiN interface. 

Figure 1-2. Effect of the surface pre-treatment with NH3 during SiN-(2) 

deposition of EFG Si without the pre-treatment (A) and with the pretreatment (B). 

36

18

6

2

0

5

10

15

20

25

30

35

40

As-grown P

diffused

A B

Bulk Lifetime (μs)

Figure 1-3 shows an identical experiment performed for the SiN-(1) film, which was 

deposited at 13.56 MHz PECVD reactor. Contrary to the SiN-(2), the SiN-(1) did not show 

the beneficial effect of the NH3 pre-treatment. This could be related to lower surface 

damage or less hydrogen implantation due to lower ion energy at this high frequency. 

1.9

6.1

8.0 8.1

0

1

2

3

4

5

6

7

8

9

10

As-grown PdiffusedA B

Bulk Lifetime (μs)

Figure 1-3. Effect of the surface pre-treatment with NH3 on SiN-(1)-induced 

defect passivation of EFG Si. 

3
In order to understand the relative importance of the release and diffusion of hydrogen 

versus retention of hydrogen for the defect passivation induced by the two SiN films, 

annealing temperatures were varied in the range of 600-900 °C in RTP. Figure 1-4 shows 

that the optimal annealing temperature for SiN-(2) was found to be ~750 °C as opposed to 

~800 °C for SiN-(1). In addition, SiN-(2) increased the bulk lifetime to ~48 μs at ~750 °C 

while SiN-(1) gave a bulk lifetime of ~25 μs at ~800 °C, indicating that SiN-(2) is more 

effective in passivating defects in EFG Si. These results reveal that the optimal cell process 

sequence should be quite different for the two SiN films in order to maximize the hydrogen 

passivation of defects. 

1.1.2 Impact of As-Grown Bulk Lifetime on SiN-(2)-Induced Defect Passivation of 

Screen-Printed EFG Silicon Solar Cells 

In this section, screen-printed solar cells were fabricated on the SiN-(2)-coated EFG Si 

wafers with as-grown bulk lifetime in the range of 1-5 μs to study the impact of as-grown 

bulk lifetime on SiN-(2)-induced defect passivation. Since the two SiN films showed a 

very different degree of defect passivation as a function of the annealing temperature 

(Figure 1), the cells were fabricated using different firing schemes to maximize 

hydrogenation. The SiN-(1)-coated cells were fabricated, involving Al-BSF firing at 

850 °C and Ag front contact firing at 740 °C in belt processing. For the SiN-(2)-coated 

cells, screen-printed Al on the back and Ag grid on the front were fired simultaneously at 

700 °C for 1 s in RTP to achieve > 30 μs bulk lifetimes (Figure 1). 

Figure 1-5 shows the cell efficiency of the screen-printed EFG Si solar cells with twodifferent SiN films. It was found that the SiN-(2)-coated cell efficiencies were much less 

dependent on the as-grown bulk lifetime compared to the SiN-(1)-coated cells. For 

example, the SiN-(2)-coated cells showed an absolute efficiency variation of less than 1% 

as opposed to ~3% for the SiN-(1)-coated cells for the as-grown bulk lifetimes in the range 

Figure 1-4. Bulk lifetime enhancement in EFG Si after annealing of SiN-(1) 

and SiN-(2) in RTP. 

40

45

50

SiN-(2)

SiN-(1)

as-grown

0

5

10

15

20

25

30

35

600 650 700 750 800 850 900

Anneal Temperature (oC)

Bulk Lifetime (μs)

4
of 1-5 μs. However, the efficiency difference between the two gradually decreased as the 

as-grown bulk lifetime increased. There was very little efficiency difference for the asgrown bulk lifetime of > 5 μs, which increased to ~50 μs for both the SiN films, using the 

different process sequences. For the as-grown lifetime of ~1 μs, the SiN-(1) gave a cell 

efficiency of ~11% while the SiN-(2) resulted in ~14% efficiency. Thus the SiN-(2) is 

more effective in improving poor quality as-grown EFG Si wafers compared to the SiN-

(1). 

1.2 Process Design, Fabrication, and Analysis of High-Efficiency ScreenPrinted EFG Si Solar Cells 

A co-firing process in a conveyor belt furnace is commonly used in the PV industry to 

manufacture low-cost Si solar cells, where PECVD SiN AR coating and screen-printed (or 

pad-printed) Al-BSF and contacts are fired simultaneously. However, a short-and lowtemperature belt co-firing could sacrifice the quality of Al-BSF in favor of contact quality. 

EFG Si solar cells (2×2 cm2

) were fabricated using the two SiN films according to the 

four process sequences shown in Figure 1-6. The complete solar cells were analyzed to 

quantify the effect of each process sequence on bulk lifetime and solar cell parameters. 

Three separate solar cell runs were made for each process to account for the statistical 

variation in the as-grown property of EFG Si and validate the reproducibility of the cell 

parameters. 

Figure 1-5. Efficiency of the screen-printed EFG Si cells with twodifferent SiN films as a function of as-grown bulk lifetime. 

9

10

11

12

13

0.0 1.0 2.0 3.0 4.0 5.0 6.0

As-Grown Lifetime (μs)

 Cell Efficiency (%)

14

15

SiN (2)

SiN (1)

5
P diffusion

PECVD SiN AR coating

1.2.1 Development of Optimal Process Sequence for EFG Si Solar Cells with HighFrequency PECVD SiN 

Table 1-2 summarizes the I-V results of the SiN-(1)-coated EFG Si cells fabricated 

using the four different process sequences outlined in Figure 1-6. The table shows the 

average, highest, and lowest efficiencies for about 55 EFG Si cells fabricated for each 

process scheme. 

Table 1-2. I-V results of the screen-printed SiN-(1)-coated EFG Si solar cells for four 

different process schemes. 

Process Efficiency (%) 

scheme 

Voc 

(mV) 

Jsc 

(mA/cm2

) FF 

Average High Low 

B-C 534.1 26.8 0.740 10.6 12.4 9.5 

R-C 547.3 29.1 0.741 11.8 12.8 11.1 

B/B 553.6 30.1 0.743 12.4 13.1 11.1 

B/R 572.9 32.1 0.749 13.8 14.6 12.9 

It was found that the co-firing schemes in the belt (Scheme B-C) and RTP (Scheme R-C) 

were less effective than the two-step firing schemes in the belt (Scheme B/B). Scheme B/R 

resulted in an average efficiency of 13.8%, which is ~1.4% higher in absolute efficiency 

than the scheme B/B, due to significantly higher Voc and Jsc. This indicates that the B/R 

Al screen-printing on the back

Ag grid screen-printing on the front

Belt contact firing

(740 oC/30 s)

I-V measurement

RTP contact firing

(700 oC/1 s)

Anneal of SiN and Al

in belt processing 

(~850 oC)

P diffusion

PECVD SiN AR coating

Al screen-printing on the back

Ag grid screen-printing on the front

Belt contact firing

(740 oC/30 s)

RTP contact firing

(700 oC/1 s)

Anneal of SiN and Al

in belt processing 

(~850 oC)

I-V measurement

B-C B/B B/R R-C

Figure 1-6. A design of four process sequences examined in this study 

for achieving effective defect passivation, Al-BSF, and contacts in EFG 

Si cells. 

6
process scheme not only improves the Al-BSF quality but also enhances the defect 

passivation in the EFG Si cells. 

Bulk lifetime is not so important for the FZ Si cells because the minority carrier 

diffusion length is much greater than the substrate thickness, both before and after cell 

processing. However, in low-cost EFG Si, the diffusion length in the as-grown state is 

generally much smaller than the substrate thickness, and it improves during the cell 

processing. As a result, both Voc and Jsc become very sensitive to process-induced bulk 

lifetime enhancement. Therefore, bulk lifetime measurements were performed on selected 

EFG Si samples before and after the cell processing to evaluate the effect of each process 

scheme on final bulk lifetime. In order to decouple the impact of contact firing induced 

hydrogenation, bulk lifetime was also measured after the 850 °C firing in the belt. This 

sample or process will be referred to as “B”. Figure 1-7 shows the final bulk lifetime for 

the five different process schemes along with the as-grown bulk lifetime, which was 

somewhat low (1-2 μs) in this experiment. The figure shows the average values along with 

the standard deviation for 24 measurements. 

1.6

6.8 6.6

13.7

11.1

14.0

0

2

4

6

8

10

12

14

16

18

20

AsgrownB-C R-C B B/B B/R

Firing Scheme

Bulk Lifetime (μs)

Figure 1-7. Bulk lifetime measurement results of the SiN-(1)-coated EFG 

Si for five different process schemes. 

The average bulk lifetime went up from 1-2 μs to only 6.8 μs and 6.6 μs after the cofiring schemes B-C and R-C, respectively. This indicates that the co-firing of the SiN-(1) 

and Al at 700-740 °C is not very effective for the bulk lifetime enhancement in EFG Si. 

After firing the SiN-(1) and the Al simultaneously at 850 °C in the belt furnace, without a 

subsequent contact firing (Scheme B), the average bulk lifetime increased to 13.7 μs, 

indicating that the 850 °C firing in the belt furnace is much more effective in passivating 

defects than the co-firing at 700-740 °C either in the belt or RTP. A subsequent contact 

firing in the belt (Scheme B/B) was found to reduce the hydrogenated bulk lifetime to ~11 

μs while the RTP contact firing (Scheme B/R) retained the hydrogenated bulk lifetime at 

~14 μs. This difference is attributed to a higher retention probability of hydrogen at defects 

during the shorter and faster RTP contact firing at ~700 °C for ~1 s compared to the longer 

and slower belt contact firing at ~740 °C for 45 s, after the hydrogenation by the process B. 

7
We found that, in very few cases, the average bulk lifetime did not correspond to the 

cell efficiency (Table 1-2 and Figure 1-7) because of the spatial non-uniformity of bulk 

lifetime in EFG Si. This is because very low lifetime regions, if present, can have greater 

influence on the cell performance. Since the average bulk lifetime values for the co-firing 

schemes B-C and R-C correspond to much smaller diffusion length (L < 200 μm) 

compared to the ~300 μm thick EFG Si substrates, the Al-BSF quality is not expected to 

have a significant effect on the cell efficiency. This is also supported by PC1D simulations, 

which showed a calculated efficiency difference of less than 0.5% for 300 μm-thick 

devices with diffusion length of less than 200 μm and BSRV values in the range of 102 to 

104 cm/s. Thus the higher quality Al-BSF produced by the two-step firing scheme B/R, 

compared to the B/B and co-firing schemes, should have negligible effect on these EFG Si 

cells because diffusion length to the substrate thickness (L/W) ratio is much less than one. 

However, the EFG Si cell data in Table 1-2 show that the two-step firing scheme involving 

the RTP contact firing (Scheme B/R) produced 3.2% and 1.4% increase in absolute 

efficiency compared to the B-C and B/B processes, respectively. Since improved BSF due 

to the RTP contact firing cannot fully account for such huge increase in EFG Si cell 

efficiency, the RTP contact firing-induced improvement has to be attributed to the bulk 

lifetime enhancement. This is directly supported by the measured bulk lifetime data in 

Figure 1-7. The greater retention of the hydrogen at the defects during the RTP contact 

firing is the result of shorter firing time as well as the faster cooling rate (> 40 °C/s) 

compared to the slow cooling (< 10 °C/s) in the belt contact firing used in this study. 

Higher cooling rates help in retaining more hydrogen at the defects, resulting in a more 

effective defect passivation and higher bulk lifetime. Therefore, we conclude that the RTP 

contact firing improves the Al-BSF quality (or BSRV) due to high ramp-up rate and 

enhances bulk lifetime due to fast cooling. High-quality FZ Si cells, where L/W is much 

greater than one, primarily benefit from the improved BSRV; while the low-cost EFG Si 

cells, where L/W is smaller than one, benefit mostly from the improved bulk lifetime due to 

higher retention of hydrogen at defects. 

It is important to note that this study was conducted using SiN-(1) film, which is not as 

effective as SiN-(2) when the as-grown lifetime value is low (1-2 μs), as shown in Section 

1. Since the as-grown lifetime of EFG Si used in this study was less than 2 μs, the 

processed lifetime values with the SiN-(1) were well below 50 μs and the cell efficiencies 

were less than 15%. In Figure 1-5, we were able to achieve 14.5-15% efficient cells with 

SiN-(1) because the as-grown bulk lifetime was ~5 μs, as opposed to less than 2 μs. Thus 

SiN-(1) is not capable of raising the bulk lifetime of average EFG Si to over 50 μs using 

low-cost manufacturable technologies. Since Section 6.9 showed that SiN-(2) is more 

effective in hydrogenating low-lifetime EFG Si, the next section will focus on the 

fabrication and analysis of SiN-(2)-coated EFG Si cells to achieve higher efficiencies. 

8
1.2.2 Development of Optimal Process Sequence for EFG Silicon Solar Cells with LowFrequency PECVD SiN 

The previous section showed that the best process sequence for SiN-(1)-coated EFG Si 

cells involved two-step firing (Scheme B/R). This process produced good BSF and 

contacts. However, hydrogenation was not sufficient to raise bulk lifetime to ~50 μs 

because of the fairly low as-grown lifetime of 1-2 μs. This resulted in average cell 

efficiencies below 14.0% with a maximum of ~14.5%. Therefore, in this section, EFG Si 

cells (2×2 cm2) were fabricated using SiN-(2) using the same four process schemes (Figure 

1-6). The as-grown bulk lifetime of the EFG Si used in this study was in the range of 1-2 

μs. 

 Table 1-3 shows the I-V results of the EFG Si cells with SiN-(2) for the four different 

process schemes. The two-step firing scheme B/R, which was found to be best for SiN-(1), 

produced an average efficiency of 13.8% with a maximum efficiency of 14.9%. Contrary 

to SiN-(1) coating, the RTP co-firing (Scheme R-C) was found to be best for the SiN-(2)-

coated EFG cells, resulting in an average efficiency of 14.3% with a maximum efficiency 

of 15% due to appreciably higher Jsc and Voc. Even the belt co-firing (Scheme B-C) was 

found to be more effective than the two-step B/B process. Thus the co-firing schemes seem 

to be better for SiN-(2) as opposed to two-step firing schemes for SiN-(1) for these four 

process schemes. 

Table 1-3. I-V results of the screen-printed SiN-(2)-coated EFG Si solar cells for four 

different process schemes 

Process Efficiency (%) 

Scheme 

Average 

Voc (mV) 

Average 

Jsc (mA/cm2

) 

Average 

FF Average High Low 

B-C 571.0 31.1 0.760 13.5 13.8 12.9 

R-C 588.6 32.0 0.761 14.3 15.0 13.4 

B/B 572.1 30.5 0.742 13.0 14.0 11.3 

B/R 582.9 31.6 0.752 13.8 14.9 13.0 

It is important to recognize that the overall cell efficiency of the SiN-(2)-coated EFG Si 

cells is appreciably higher than the SiN-(1)-coated cells for all the corresponding process 

schemes. These results are supported by the bulk lifetime measurements in Figure 1-8. The 

R-C process increased the average bulk lifetime from 1-2 μs to ~35 μs but is still less than 

50 μs. Therefore, in the next section, we tried to optimize the hydrogenation temperature 

for SiN-(2) using the R-C process to achieve bulk lifetime over 50 μs. 

9
20.5

12.9

34.7

18.2

1.7

0

5

10

15

20

25

30

35

40

45

50

AsgrownB-C R-C B/B B/R

Firing Scheme

Bulk Lifetime (μs)

Figure 1-8. Bulk lifetime measurement results of the SiN-(2)-coated EFG 

Si for four different process schemes. 

1.2.3 Optimization of Hydrogenation Temperature in RTP for Low-Frequency PECVD 

SiN to Achieve High-Efficiency EFG Si Cells with Photolithography Contacts 

Since the optimum hydrogenation temperature of ~850 °C for the SiN-(1) in the belt 

furnace, the resulting optimum process sequence involved two-step firing (850 oC/2 min in 

the belt furnace and 700 °C/1 s in RTP), as discussed earlier. This sequence produced good 

contacts, BSF, and fairly effective hydrogenation. In the previous section, for the SiN-(2) 

we found that a relatively low temperature (700 °C) and short time (~1 s) RTP firing with 

fast cooling rate (> 40 °C/s) was very effective in producing good EFG Si cells (~15% 

efficiency). However, the firing condition may not be optimal for hydrogenation because 

the 700 o

C/1 s firing condition was tailored to maintain good FF or screen-printed Ag 

contacts. Therefore, in this section, a study was conducted to determine the precise 

optimum hydrogenation temperature for the SiN-(2) in RTP and utilize it to achieve highefficiency EFG cells first by using photolithography contacts, where we don’t have to 

worry about the post-hydrogenation heat treatment because they are formed at the end. 

1.2.4 Optimization of Hydrogenation Temperature for SiN-(2) in RTP 

 In this study, 10×10 cm2

 EFG Si wafers from the same growth were subjected to the P 

diffusion followed by the PECVD SiN-(2) coating on top of the emitter. After the screenprinting of Al on the back, the samples were fired in the RTP system at temperatures in the 

range of 600-850 °C for 60 s with two different cooling rates. To evaluate the degree of the 

SiN-(2)-induced hydrogenation, the fired samples were etched down to bare Si and 

immersed in the iodine-methanol solution for surface passivation for the bulk lifetime 

10
measurement. The bulk lifetime was measured on ten different locations on each wafer and 

two wafers were used for each firing condition to account for nonuniformity. 

Figure 1-9 shows the average bulk lifetime of the samples after the RTP firing of SiN-

(2) and Al with two different cooling rates. The fast cooling rate (~40 °C/s) was ~40 times 

greater than the slow cooling rate. It was found that the slow cooling condition showed 

maximum bulk lifetime improvement at ~650 °C with the average lifetime of ~36 μs, as 

opposed to ~48 μs with a maximum of 74 μs for the ~750 °C anneal with fast cooling. This 

is the result of the competition between the release of atomic hydrogen from SiN, its 

diffusion through the bulk, and retention of hydrogen at the defects. These results show 

enhanced hydrogen passivation of defects at faster cooling rate due to higher retention of 

hydrogen at the defects. In addition, the 700 °C/60 s annealing increased the average bulk 

lifetime to ~40 μs, but the 850 °C/60 s annealing resulted in an average lifetime of only 

~17 μs. This result explains why the two-step firing, involving the 850 °C hydrogenation in 

the previous section, was not very effective as it could be for SiN-(2). Furthermore, it also 

shows that 850 °C RTP firing is not appropriate for the SiN-(2)-induced hydrogenation. 

Figure 1-9 reveals that the 700 °C/1 s co-firing, which gave ~15% efficiency in the 

previous section, is sub-optimal for SiN-(2)-induced hydrogenation. In addition, the 750°C 

RTP firing, which is the optimum for the SiN-(2)-induced hydrogenation, is not good for 

screen-printed contacts because it gave low FF. These conflicts show the challenge of 

developing an optimal process sequence for EFG Si cells. Therefore, in the next section, 

EFG Si cells were fabricated with photolithography contacts first, using the optimum 

hydrogenation condition of ~750 °C/60 s in RTP. 

10

15

20

25

30

35

40

45

50

55

550 600 650 700 750 800 850 900

Temperature (oC)

Bulk Lifetime (μs)

slow cooling

fast cooling

Figure 1-9. The average bulk lifetime along with standard deviation of EFG 

Si after the simultaneous firing of SiN-(2)/Al at temperatures of 600-850 °C 

in RTP. The firing time was fixed at 60 s. 

11
1.3 Fabrication and Analysis of 15.8% Efficient EFG Si Solar Cell with 

Photolithography Contacts 

In the previous section (Figure 1-9), we established that the optimal hydrogenation 

condition for SiN-(2) involved 750 °C /60 s firing in RTP, which could give the average 

bulk lifetime of ~50 μs required for ~16% efficient EFG Si cells with BSRV of 500-1000 

cm/s. Unfortunately, the ~750 °C firing in the RTP is not suitable for screen-printed Ag 

grid contacts because it drives the metal too deep, causing junction shunting and low FF. 

Therefore, in this section, we decided to form evaporated metal contacts using lift-off 

photolithography after the hydrogenation of the cells at ~750 °C, Figure 1-10. This should 

result in the best hydrogenation and good contacts. 

P diffusion

PECVD SiN deposition

Al printing on the back

Ag grid printing

RTP co-firing

(700 oC/1 s) Metal evaporation

Lift-off

I-V measurement

RTP firing

(700-850 oC/60 s)

P diffusion

PECVD SiN deposition

Al printing on the back

Ag grid printing

RTP co-firing

(700 oC/1 s) Metal evaporation

Lift-off

I-V measurement

RTP firing

(700-850 oC/60 s)

Figure 1-10. Process sequence and conditions used for fabricating the recordhigh efficiency EFG Si cells with the SiN-(2) and RTP firing. The left-hand 

side shows an industry screen-printed cell processing. 

Table 1-4 shows the average efficiencies of the EFG cells (2×2 cm2) as a function of 

the peak firing temperature in the range of 700-850 °C. As expected from the lifetime 

results in Figure 1-9, the ~750 °C firing was found to be best for maximum defect 

passivation, resulting in an average efficiency of 15.7% with a maximum of 15.8%. The 

efficiency drops on either side of ~750 °C. The reason for this optimum is that less 

hydrogen is released and diffused at temperatures below 750 °C while more hydrogen is 

dissociated from the defects at temperatures above 750 °C. This is supported by the bulk 

lifetime in Figure 1-9 and Jsc and Voc values in Table 1-4. LBIC measurements were 

performed to map the electrical activity of defects over the cell area. In this technique, a 

He-Ne laser with wavelength of 905 nm is used to generate carriers, and the generated 

carriers are collected in short circuit mode to assess the recombination. Dark regions reflect 

less current collected or higher recombination. LBIC measurements in Figure 1-11 clearly 

show less recombination (more lighted regions) and higher spectral response (A/W) for the 

15.8% efficient cell fired at ~750 °C. 

12
Table 1-4. The result of I-V measurements on the EFG Si cells with photolithography 

contacts, in which the SiN-(2)/Al firing temperature was varied in the range of 700-850 °C 

in RTP. 

Firing Temp. 

(

o

C) 

Voc 

(mV) 

Jsc 

(mA/cm2

) FF Eff. (%) 

Rs 

(Ω-cm2

) 

Rsh 

(Ω-cm2) 

700 584 32.82 0.771 14.8 0.89 44193 

750 603 33.19 0.785 15.7 0.60 50778 

800 590 32.91 0.772 14.9 0.79 41321 

850 567 30.26 0.764 13.1 0.64 42794 

It is important to point out that once the SiN film was fired at ~750 °C it became very 

difficult to etch for a good lift-off photolithography. This caused appreciable amount of 

undercutting, resulting in somewhat wider openings than the photolithography grid design. 

In fact, the loss of current due to wider grid and FF of only ≤ 0.77 for these 

photolithography cells suggests that manufacturable screen-printed cells, which have wider 

grid and lower FF, with this hydrogenation condition could also give ~16% efficient cells. 

Therefore, in the next section, we developed a new process sequence using two-step firing 

in RTP to retain maximum hydrogenation while maintaining good screen-printed contacts. 

13
Firing temperature: 700 °C 

Efficiency: 15.2% 

Average A/W: 0.552 

Firing temperature: 750 °C 

Efficiency: 15.8% 

0.66 Average A/W: 0.570 

0.64

0.62

0.60

0.58

0.56

0.54

0.52

0.50

0.48

0.46

0.44

0.42

0.40

0.38

0.36

Amps/Watt

0.66

0.64

0.62

0.60

0.58

0.56

0.54

0.52

0.50

0.48

0.46

0.44

0.42

0.40

0.38

0.36

Amps/Watt

0.66

0.64

0.62

0.60

0.58

0.56

0.54

0.52

0.50

0.48

0.46

0.44

0.42

0.40

0.38

0.36

0.66

0.64

0.62

0.60

0.58

0.56

0.54

0.52

0.50

0.48

0.46

0.44

0.42

0.40

0.38

0.36

Amps/Watt

Firing temperature: 800 °C 

Efficiency: 15.0% 

Average A/W: 0.545 

Figure 1-11. LBIC response of EFG Si cells fired at 700 °C, 750 °C, and 800 °C for 60 s 

in RTP. The values in the legend represent photo-response in A/W. 

14
1.4 Fabrication and Analysis of Record-High Efficiency Screen-Printed EFG 

Silicon Solar Cell 

Since low-cost screen-printed contacts are much more desirable and cost effective than 

the photolithography contacts, a new process sequence was established in this section that 

not only retains maximum defect passivation but also produces high-quality screen-printed 

contacts. Since the optimal hydrogenation condition (750 °C/60 s firing in RTP) is not 

appropriate for screen-printed contacts, we decided to use two-step firing: one for optimal 

hydrogenation and the second for good screen-printed contacts without loosing the 

hydrogen at the defects. Since post-hydrogenation anneal at high temperatures can cause 

dehydrogenation of defects, in this study, first we examined the extent of rapid cooling 

required for retaining the hydrogen at the defects. 

EFG Si wafers were subjected to P diffusion, SiN-(2) deposition, and Al screenprinting, followed by an RTP annealing at ~750 °C for 60 s for maximum hydrogenation 

of the samples. After the hydrogenation, the samples were re-annealed in RTP to simulate 

contact firing at 700 °C for 1 s with four different cooling rates. Bulk lifetime 

measurements were performed to assess the impact of cooling rate associated with the RTP 

contact firing. This is important because hydrogen dissociates from the defects in EFG Si 

during post-hydrogenation annealing. 

Figure 1-12 shows the relative change in bulk lifetime of hydrogenated EFG Si due to 

the RTP contact firing. As expected, the hydrogenated bulk lifetime decreased sharply for 

slow cooling rates. The slowest cooling rate (~1 °C/s) caused more than 50% reduction in 

bulk lifetime. This is attributed to reactivation of defects, resulting from evolution of the 

hydrogen from defects. At the cooling rate of ~30 °C/s, the reduction in the bulk lifetime 

was only ~3%. The hydrogenated bulk lifetime was completely preserved at the cooling 

rates of ≥ 40 °C/s. 

40

50

60

70

80

90

100

1 10 30 40

Cooling Rate (oC/s)

Relative Bulk Lifetime (%)

Figure 1-12. Relative bulk lifetime change of hydrogenated EFG Si 

due to RTP contact firing at ~700 °C for ~1 s with four different 

cooling rates. 

15
In the previous sections, optimum hydrogenation condition for the SiN-(2) was found 

to be ~750 °C/60 s with a cooling rate of ≥ 40 °C/s, and the 700 oC/1 s RTP firing with ≥

40 °C/s cooling rate could preserve hydrogenated bulk lifetime. In addition, we found that 

it could also produce good screen-printed contacts. This led to EFG Si cell fabrication with 

two-step RTP firing, involving ~750 oC/60 s for defect passivation using the SiN-(2) 

followed by ~700 °C/1 s for screen-printed contacts. Fast cooling rate of 40 °C/s was 

implemented for both steps. 

Figure 1-13 shows a process sequence designed to take full advantage of the two 

optimized RTP firing steps for defect passivation and screen-printed contacts. Fabrication 

of EFG Si cells (2×2 cm2) involved manufacturable P diffusion, PECVD SiN-(2) 

deposition, screen-printed metals, and two-step RTP firing. In this experiment, the RTP 

hydrogenation temperature was varied in the range of 730-780 °C to reestablish and 

support maximum defect passivation temperature of 750 °C for complete solar cells. As 

many as 61 EFG Si cells in three separate runs were fabricated to account for the statistical 

variation in property of EFG Si. The as-grown bulk lifetime was 1-2 μs, which represents 

average-quality EFG Si wafers. 

P diffusion (930 oC/6 min) in belt processing 

PECVD SiN-(2) deposition

Al screen-printing on the back

SiN-(2)/Al firing (730-780 oC/60 s) in RTP

Ag grid screen-printing on the front

Contact firing (700 oC/1 s) in RTP

I-V measurement

P diffusion (930 oC/6 min) in belt processing 

PECVD SiN-(2) deposition

Al screen-printing on the back

SiN-(2)/Al firing (730-780 oC/60 s) in RTP

Ag grid screen-printing on the front

Contact firing (700 oC/1 s) in RTP

I-V measurement

Figure 1-13. A process sequence for record high efficiency screenprinted EFG Si cells, involving two-step RTP for defect passivation and 

good contacts. 

Tables 1-5A and 1-5B summarize the results of the screen-printed EFG Si cells as a 

function of the RTP hydrogenation temperature. Both average (Table 1-5A) and best 

(Table 1-5B) cell parameters for each hydrogenation temperature are shown. Note that 

~750 °C hydrogenation in RTP was again found to be the optimum for defect passivation, 

resulting in average cell efficiency of 15% with a maximum of 15.9% for about 60 cells. 

Figure 1-14 shows that 15.9% represents the highest efficiency for screen-printed (or padprinted) Si ribbon solar cells to date. Figure 1-15 shows the lighted I-V data for this recordhigh efficiency (15.9%) EFG Si cell (verified by National Renewable Energy Laboratory). 

This cell has excellent Voc of 615 mV, Jsc of 33.7 mA/cm2

, and FF of 0.769. This is the 

result of high bulk lifetime of > 50 μs, fairly good BSF, and high-quality screen-printed 

16
contacts. Tables 1-5-A and 1-5-B show that for hydrogenation temperatures above 750 °C, 

the cell efficiency decreases rapidly because of reduced retention of hydrogen at defects. 

This is reflected in decrease of the Jsc and Voc values, despite better Al-BSF at higher 

temperature. It is noteworthy that the FFs are in the range of 0.761-0.783 because of very 

good contacts with low Rs ( 0.6-0.8 Ω-cm2

) and high Rsh (> 5×104 Ω-cm

2

). This 

demonstrates that the short RTP contact firing at 700 °C is beneficial to both 

hydrogenation and screen-printed contacts due to fast cooling and heating rates. 

Table 1-5A. The average parameters of the screen-printed EFG Si cells, in which two-step 

RTP firing was involved for defect passivation and contacts. 

Hydrogenation 

Temp. (°C) 

Voc 

(mV) 

Jsc 

(mA/cm2

) FF Efficiency (%) 

730 591 31.6 0.771 14.4 

740 595 31.8 0.776 14.7 

750 601 32.2 0.773 15.0 

760 596 31.7 0.768 14.5 

770 590 31.2 0.763 14.1 

780 583 30.9 0.765 13.8 

Table 1-5B. The best parameters of the screen-printed EFG Si cells, in which two-step RTP 

firing was involved for defect passivation and contacts. 

Hydrogenation 

Temp. (°C) 

Voc 

(mV) 

Jsc 

(mA/cm2

) FF Efficiency (%) 

730 605 32.6 0.771 15.2 

740 607 32.8 0.773 15.4 

750 615 33.7 0.769 15.9 

760 610 33.1 0.770 15.5 

770 600 32.2 0.781 15.1 

780 598 31.6 0.771 14.6 

17
15.0

15.3 [31]

15.9

14.5 [45]

12.5 [43]

12.0 [32]

15.4 [21]

14.7 [46]

14.2 [18]

12.0

12.5

13.0

13.5

14.0

14.5

15.0

15.5

16.0

16.5

1980 1985 1990 1995 2000 2005

Year

Cell Efficiency (%)

EFG

String Ribbon This study

Web

Figure 1-14. Efficiencies of screen-printed Si ribbon solar cells, 

involving the highest efficiency of 15.9% achieved in this thesis. 

0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

0.1 0.2 0.3 0.4 0.5 0.6 0.7

Voltage (V)

Current (A)

Voc (mV) Jsc (mA/cm2

) FF Eff. (%)

615 33.7 0.769 15.9

Figure 1-15. I-V curve for the 15.9% efficient screen-printed EFG Si 

solar cell, measured by National Renewable Energy Laboratory. 

18
Figure 1-16 shows the efficiency distribution of the screen-printed EFG Si cells that 

were hydrogenated at the optimum temperature of ~750 °C. Efficiencies of about 60 cells 

were in the range of 14-16% with the most number of cells at ~15.2. The wide spread in 

efficiency is attributed to non-uniform distribution of impurities and defects in EFG Si. 

Figure 1-17 shows LBIC measurement on nine 2×2 cm2

 cells made on one 10×10 cm2 EFG 

Si wafer. LBIC measurements show considerable nonuniformity of electrically active 

defects over the full wafer, even after the optimum hydrogenation. Although the efficiency 

of most cells on this wafer is ≥ 15%, Figure 1-17 shows a good correlation between the cell 

efficiency and the LBIC map. 

0

2

4

6

8

10

12

14

16

14.2 14.4 14.6 14.8 15.0 15.2 15.4 15.6 15.8 16.0

Cell Efficiency (%)

Number of Cells

Figure 1-16. An efficiency distribution of 61 screen-printed 

EFG Si cells involving two-step RTP firing scheme. 

19
Cell 

No. 

Voc 

(mV) 

Jsc 

(mA/cm2

) FF Eff. (%) 

LBIC response 

(A/W) 

1 599 32.6 0.759 14.9 0.532 

2 600 32.7 0.758 14.9 0.541 

3 614 33.2 0.764 15.7 0.593 

4 604 32.8 0.762 15.0 0.554 

5 609 32.5 0.759 15.1 0.580 

6 617 33.3 0.767 15.7 0.593 

7 573 31.3 0.762 13.7 0.538 

8 607 33.2 0.758 15.3 0.579 

9 616 33.5 0.769 15.9 0.601 

20
1.5 Conclusion 

Hydrogenation due to PECVD SiN deposited with a frequency of ~150 KHz at ~400 °C 

was found to be superior to the SiN deposited with a frequency of 13.56 MHz at 300 °C, 

despite the fact that the high-frequency SiN film contained more atomic hydrogen in the 

form of Si-H and N-H bonds. This suggested an additional source of hydrogen for the lowfrequency SiN film, which was attributed to NH3 pre-treatment during the low-frequency 

SiN deposition. The low-frequency SiN film was found to be more effective in passivating 

defects in EFG Si with low as-grown bulk lifetimes (< 3 μs). The cell efficiency difference 

between the two SiN films decreased as the as-grown bulk lifetime increased to ~ 5 μs. 

The low-frequency SiN film showed an optimum hydrogenation temperature of ~750 °C 

while the high-frequency SiN film showed the optimum at ~800 °C in RTP. A ~2 μs asgrown lifetime increased to ~25 μs after 850 °C firing and ~50 μs after ~750 °C firing of 

the high- and low-frequency SiN films, respectively. 

A complete process sequence was developed through fundamental understanding to 

achieve record-high efficiency EFG Si cells. For SiN-(1)-coated EFG Si cells, a two-step 

process scheme (~850 °C/2 min in the belt + ~700 °C/1 s in RTP) was found to be optimal 

for defect passivation and screen-printed contacts, resulting in an average efficiency of 

~13.5% with a maximum of 14.6%. RTP contact firing with fast ramp-up and cooling gave 

~1.4% increase in absolute cell efficiency compared to the slow belt firing used in 

production. This improvement was largely due to bulk lifetime enhancement associated 

with rapid cooling in RTP, which enhances the retention of hydrogen at the defects in EFG 

Si cells. 

While a two-step firing, involving hydrogenation at 850 °C, was found to be best for 

EFG Si cells with SiN-(1), it was not appropriate for SiN-(2). This is because optimal 

hydrogenation temperature was found to be ~750 °C for SiN-(2)-induced defect 

passivation, compared to ~850 °C for SiN-(1). Since ~750 °C was too intense for screenprinted contacts, a two-step firing scheme was developed by optimizing the firing of 

screen-printed contacts after the hydrogenation. It was found that the RTP contact firing 

(~700 °C/1 s) with fast cooling rate (≥ 40 °C/s) could preserve the hydrogenated bulk 

lifetime due to higher retention of hydrogen a the defects. The two-step RTP processes, 

involving defect passivation at ~750 o

C/60 s and screen-printed contact firing at 

~700 °C/1 s, produced the record-high cell efficiency of 15.9% on Si ribbon to date. 

21
2. RTP-Enhanced Hydrogen Passivation in Ribbon 

Multicrystalline Silicon Studied by Scanning RoomTemperature Photoluminescence Spectroscopy 

2.1 Introduction 

Impressive growth in the photovoltaic market over the past decade has been 

predominantly driven by advances in crystalline silicon technology. Multi-crystalline 

silicon (mc-Si), which is produced by many competing techniques, can meet low-cost 

production, high solar cells throughput and high efficiency requirements. The importance 

of the defect passivation or elimination in low-cost mc-Si wafers is recognized as a key 

requirement for achieving high efficiency solar cells. Solar cell production employs a set of 

process steps, which realize improvement in material’s quality. These manufacturing steps 

include (a) phosphorous diffusion, (b) hydrogenation from the amorphous SiNx:H layer, 

and (c) Al back contact firing. It was established recently that combined effect of 

simultaneous (b) + (c) steps using rapid thermal processing (RTP) technique provides 

additional benefits to solar cell efficiency through the hydrogen mediated defect reduction 

[2-3]. This new hydrogen passivation approach tentatively accounted for increased 

hydrogen incorporation in defect-enriched regions of mc-Si wafers and solar cells. In this 

paper, we report on scanning room-temperature photoluminescence (PL) spectroscopy 

study of the simultaneous rapid thermal annealing of Al on the back and SiNx on the front 

(RTP-Al/SiNx) hydrogenation mechanism. 

2.2 Solar Cell Fabrication And Photoluminescence Analysis Procedures 

For the PL excitation, the 800nm pulse AlGaAs laser diode with 140 mW peak power 

was used. PL mapping set-up and procedure are published elsewhere [2-4]. Materials used 

in this study were boron doped ~300 μm thick ribbon mc-Si wafers grown by Edgedefined Film-fed Growth (EFG) method. A set of commercial EFG wafers was prepared 

and processed as specified in Table 1. Each processing step was followed with PL mapping 

measurement at two principal luminescence bands as described below. 

The PL spectrum of mc-Si wafers at room temperature is composed of two 

luminescence bands. First band (IBB) shows a spectral maximum at 1.09eV and 

corresponds to band-to-band recombination in silicon. This band was characterized in a 

broad temperature range up to 750 K and attributed to direct recombination of free and 

bound to exciton electron-hole pairs [2-5]. The second “defect” band at 0.8 eV (IDef) is 

observed only at selected regions on the wafer, and corresponds to deep luminescent 

defects located around the dislocation lines in mc-Si [2-6]. This defect band is a finger 

print of the low lifetime regions on wafers and caused by dislocations contaminated by 

impurity precipitates [2-4]. PL maps with reasonable spatial resolution can be used to 

deduce (a) the distribution of the effective minority carrier lifetime (τeff) [2-7], (b) the 

lifetime variation after a solar cell processing, and (c) the maps of the defect centers 

causing 0.8eV PL band. The τeff is composed of radiative, non-radiative, and surface 

recombinations and in general, dominated by non-radiative recombination in Si. In a 

22
simple case neglecting a diffusion of the photo-generated minority carriers, the IBB and IDef

can be expressed as following [2-4]: 

 IBB = C1 × G × τeff / τrad (1-2) 

 IDef = C2 × G × τeff / τSRH (2-2) 

where G is the electron-hole pair’s photo-generation rate, τrad is the radiative 

recombination lifetime, τSRH = (NDef vth σn)

-1 is the Shockley–Read–Hall radiative lifetime, 

NDef is the radiative defect concentration, vth and σn are the electron thermal velocity and 

capture cross section, C1 and C2 are temperature dependent Si constants. A new parameter, 

R = IDef / IBB, can be introduced, to measure the point-by-point ratio of two PL intensities. 

Following Eqs. (2-1) and (2-2) it is clear that 

 R = IDef / IBB = const × NDef (3) 

i.e. R-parameter allows directly accessing the PL defect concentration and their spatial 

distribution with the PL mapping procedure. Notice that the R-value is independent on 

other recombination defects in the bulk and at the surface, which contribute to τeff. 

Table 2-1. Flow-chart of processing steps and PL results in EFG wafers. Average across the 

wafer values of the band-to-band PL (IBB), defect band PL (IDef) and R-parameter (R= IDef/ 

IBB) on sample #2 before and after each processing step are summarized. 

Sample stage Processing IBB IDef R 

1 Initial 28.8 *) *) 

2 Hydrogenation 

SiNx deposition, Al 

back contact screenprint, RTP annealing at 

750°C for 1sec, and 

removing SiNx layer in 

50% HF solution. 

126.3 0.71 0.011 

3 De-hydrogenation 

(600°C) 

RTP annealing at 

600°C for 1sec in N2/O2

ambient 

112.5 0.61 0.013 

4 De-hydrogenation 

(700°C) 

RTP annealing at 

700°C for 1sec in N2/O2

ambient 

42.8 0.34 0.024 

*) Defect band intensity is below sensitivity limit of the PL system. 

23
2.3 Characterization of Defects in EFG Si Before and After Hydrogenation by 

Room Temperature PL Mapping 

Figs. 2-1a and 2-1b show two PL maps, IBB and IDef, respectively, collected at identical 

sample spots on 50 mm × 22 mm EFG wafer after the 750 °C RTP-Al/SiNx hydrogenation 

step. The distribution of IBB maintains a profile of the effective lifetime according to Eq. 

(2-1) as published in [2-4]. The areas with low IBB intensity correspond to high 

recombination activity of all defects (low lifetime regions), while the high IBB values 

represent “good” or high lifetime regions. It was established previously that low IBB

regions are specified also by intense IDef band, while in high IBB regions this band is 

negligible [2-4]. This reverse contrast of IBB and IDef is illustrated in Figs. 2-1a and 2-1b 

after the hydrogenation was completed. 

Hydrogenation leads to a significant increase in the band-to-band PL intensity (28.8 to 

126.3) averaged across the sample, as presented in Table 2-1 (wafer #2). This effect can be 

interpreted as effective lifetime (τeff) upgrade due to RTP-Al/SiNx processing [2-8]. In 

different samples we observed from 1.5 to 4.3 times hydrogenation-induced increase in 

average IBB values. Notice that the defect band, which at initial stage had intensity below 

the PL system detection limit also increased in intensity after the hydrogenation due to the 

Figure 2-1. Room temperature PL mapping (a) band-to-band (b) defect band PL, and (c) 

point-by-point ratio of IBB (hydrogenated)/IBB(initial) representing lifetime upgrade effect. 

The mapping size is 50 mm x 22 mm, step=0.5 mm. 

24
same effective lifetime upgrade according to Eq. 2-2. We observe that high PL and low PL 

areas on EFG wafers respond differently to the RTP-Al/SiNx processing. Specifically, high 

IBB areas prior to the hydrogenation show quite small, up to 30%, lifetime upgrade effect. 

In contrast, low IBB areas show a spectacular lifetime upgrade, up to two orders of the 

magnitude, after the hydrogenation. Evidently, the average across the sample values of IBB

increase (Table 2-1) caused by hydrogenation represents a superposition of both areas. To 

illustrate the lifetime upgrade effect over the wafer one can introduce the upgrade 

coefficient as the point-by-point ratio of two band-to-band PL intensities, before and after 

the hydrogenation: 

Kup = IBB (hydrogenated) / IBB (initial) = τeff (hydrogenated) / τeff (initial) (2−4) 

The Kup map on the wafer #2 is plotted in Fig. 2-1c. To further quantify the lifetime 

upgrade, we show in Fig.2 PL data analyzed on two different areas on sample #2 with (a) 

“low” and (b) “high” band-to-band PL. Consistent with entire mapping in Fig. 2-1, low 

band-to-band PL area (a) shows in average 20-fold lifetime increase, while high band-toband PL area (b) shows only 10% upgrade in lifetime. Notice that the y-axis on Fig. 2-2 

has a logarithmic scale. 

Figure 2-2. Lifetime up-grade and down-grade coefficients measured in areas with (a) low 

and (b) high band-to-band PL intensity; (1) corresponds to RTP-Al hydrogenation; (2) 600 

°C RTP de-hydrogenation, and (3) 700 °C RTP de-hydrogenation. 

25
Based on the above study, we conclude that RTP-Al/SiNx hydrogenation significantly 

improves “bad” regions in EFG wafers, and has a little effect on the high lifetime areas. 

To further explore the hydrogen incorporation mechanism in mc-Si, we performed two 

consecutive RTP annealings at 600 °C and 700 °C on the previously hydrogenated wafers 

after etching off the SiNx:H layer (see Table 2-1 for the processing details). It was 

demonstrated recently that the average effective lifetime decreased after this processing 

step due to hydrogen out-diffusion, and reduction of bulk lifetime can lead to significant 

loss in solar cell efficiency [2-3]. In Table 2-1, we show PL results of such a dehydrogenation experiment on wafer #2. Consistent with the lifetime data in literature, the 

RTP at 600 °C and 700 °C led to a gradual decrease in the PL intensity of both the 

luminescent bands, which is consistent with the hydrogen out-diffusion model for the 

observed lifetime down-grade [2-8]. 

Similar to the hydrogenation case, we analyzed the IBB variation across the wafers after 

each de-hydrogenation anneal step. We calculated at every mapped spot the ratio of the 

band-to-band PL intensities after the RTP with respect to the hydrogenated stage by 

introducing the down-grade coefficient 

Kdown= IBB (de-hydrogenated) / IBB (hydrogenated) 

 = τeff (de-hydrogenated) / τeff (hydrogenated) (2-5) 

In Fig. 2-2, the Kdown data points are shown for the same wafer regions as in the 

hydrogenation experiment, allowing comparing them with Kup distributions. In addition to 

the PL degradation across the entire sample, two other effects were observed. First, the low 

PL intensity region “a” showed much stronger lifetime downgrade than high PL region 

“b”. Specifically, we observed up to 50-fold decrease in the PL intensity in the area “a” 

versus only a factor of two decrease in the area “b”. Secondly, the PL intensity degradation 

due to RTP dehydrogenations anneal showed a different temperature annealing rate in “a” 

compared to “b” regions. This is evident by comparing points labeled as (2) and (3) in both 

areas. This indicates that the activation energies of the hydrogen out-diffusion process are 

quite different in both regions; lower in the low-lifetime regions and higher in the high 

lifetime regions. Moreover, the rapid decrease in IBB in area “a” is in general agreement 

with the reported lifetime degradation [2-8]. To elaborate on details of the hydrogenation 

and de-hydrogenation processes, we present in Fig. 2-3 three PL line scans measured and 

calculated at identical wafer location. Data (a) represents the line-scan for the downgrade 

coefficient after the 700 °C RTP de-hydrogenation anneal. Notice that Kdown shows up to 

10 times reduction in lifetime in the areas marked by arrows. Data points in Fig. 2-3b 

represent R-parameter calculated according to Eq. (2-3) after the hydrogenation (open 

circles) and subsequent de-hydrogenation (closed circles) at 700 °C. First, notice that a 

clear inverse correlation between Kdown and R-parameter or the radiative defect 

concentrations. Secondly, comparison of R-parameter distribution after the RTP-Al/SiNx

hydrogenation and spatial match of strongest down-grade lifetime effect and R-parameter 

peaks indicated by arrows. Subsequent 700 °C de-hydrogenation reveals that the increase 

in R-parameter or the PL defect concentration is caused by hydrogen out-diffusion. This 

26
line-scan experiment is statistically in agreement with the behavior over entire EFG wafer 

area as illustrated in Table 2-1. 

Figure 2-3. Lines scans of (a) down-grade coefficient after de-hydrogenation cased by 700 °C RTP; 

and (b) R-parameter after RTP-Al hydrogenation (open circles) and after 700 °C de-hydrogenation 

(closed circles). 

 

2.4 Conclusions 

In conclusions, scanning spectroscopic PL experiments on hydrogenated and dehydrogenated EFG mc-Si wafers identified that the major effect of the RTP-Al/SiNx anneal 

process consists of hydrogen incorporation into the regions with high recombination 

activity of defects involving contaminated dislocations. Hydrogenation passivates radiative 

and non-radiative dislocation defects in the low lifetime regions, improving the local 

minority carrier lifetime, which benefits the solar cell efficiency. 

27
2.5 References 

[2-1] S. Narasimha and A. Rohatgi, IEEE Trans. Electron Devices 45, 1776 (1998) 

[2-2] A. Rohatgi, D. S. Kim, K. Nakayashiki, V. Yelundur, and B. Rounsaville, Appl. 

Phys. Lett. 84, 145 (2004). 

[2-3] V. Yelundur, A. Rohatgi, A. Ebong, A. M. Gabor, J. I. Hanoka, and R. L. Wallace, J. 

Electron. Mater. 30, 526 (2001). 

[2-4] S.Ostapenko, I.Tarasov, J.P.Kalejs, et al, Semicond. Sci. Technol. 15 (2000) pp.840-848. 

[2-5] V. Alex, S. Finkbeiner, J. Weber, J. Appl. Phys. 79 (1996), 6943. 

[2-6] S. Ostapenko, M. Romero European Physical Journal: Applied Physics (2004), in press. 

[2-7] I.Tarasov, S.Ostapenko, V. Feifer et al, Physica B, 273-274 (1999), pp.549-552 

[2-8] K.Nakayashihi, D.S.Kim, A.Rohatgi, B.R.Bathey, Technical Digest of 14th

International Photovoltaics Science and Engineering Conference, Bangkok, Thailand, 2004 

p. 643. 

28
3. Effect of Material Inhomogeneity on the Open-Circuit Voltage 

of String Ribbon Si Solar Cells 

3.1 Introduction 

The cost of photovoltaic (PV) systems needs to decrease by a factor of 2-4 to compete 

with traditional energy sources. Cast and ribbon multicrystalline Si (mc-Si) are promising 

low-cost materials for cost-effective PV systems. Since the Si substrate accounts for about 

40% of the total cost of crystalline Si PV modules [3-1], the use of cheaper mc-Si materials 

without appreciable loss in efficiency should help in cost reduction. In general, mc-Si 

materials are cheaper than monocrystalline Si because of the simpler and less expensive 

crystal growth technique and equipment. The ribbon mc-Si growth techniques offer an 

additional cost benefit by eliminating the kerf loss associated with the mechanical sawing 

process and subsequent deep chemical etching used for saw damage removal [3-2],[3-3]. 

Table 3-1 Characteristics of PV-grade Si materials [3-2]. 

Material 

Thickness 

(μm) 

Resistivity 

(Ω-cm) 

Carbon 

(cm-3) 

Oxygen 

(cm-3) 

Dislocations 

(cm-2) 

Czochralski 200~400 1 – 3 0.5~2.5× 1017 0.1~2.0× 1018 500 

Directional 

solidification 

300 0.8 – 1.5 < 5.0× 1017 < 2.5× 1017 1.0× 105 

EFG 300 1 – 3 1018 2.4~4.0× 1017 < 1.0× 106 

String Ribbon 300 1 – 3 4.0× 1017 < 5.0× 1016 5.0× 105 

Table 3-1 lists important characteristics of these mc-Si materials along with the PVgrade monocrystalline Czochralski Si [3-2]. Unlike monocrystalline Si, cast and ribbon 

mc-Si suffer from high concentrations of impurities and crystallographic defects that are 

distributed inhomogeneously [3-2],[3-3]. These defects lead to relatively low as-grown 

carrier lifetimes, typically <30 μs in cast mc-Si wafer and <5 μs in ribbon Si, such as edgedefined film-fed growth and String Ribbon Si [3-4]. These minority carrier lifetime values 

are not sufficient to produce high-efficiency screen-printed Si solar cells. Therefore, it is 

necessary to enhance the carrier lifetime during the solar cell fabrication process in order to 

achieve high-efficiency and cost-effective screen-printed Si solar cells. P diffusion-induced 

impurity gettering and the SiNx-induced hydrogen passivation of defects are routinely used 

as part of the solar cell fabrication sequence to enhance the carrier lifetime in these 

defective Si materials [3-4]-[3-9]. 

29
Inhomogeneously distributed electrically active defects are frequently found to be 

present in the mc-Si materials, even after effective P diffusion gettering and the SiNxinduced hydrogen passivation of defects. Therefore, understanding and assessing the 

impact of the inhomogeneous distribution of defects on solar cell performance has become 

an area of active investigation. For example, Sopori [3-10] made an attempt to quantify the 

influence of material inhomogeneities on solar cell performance by developing a 

methodology and technique to extract an effective minority carrier diffusion length (Leff) 

from the lateral and vertical distribution of defects to assess their impact on open-circuit 

voltage (VOC). Mijnarends et al. [3-11] performed numerical simulations to investigate the 

lateral variations in material quality on solar cell characteristics and found that the wider 

distribution of diffusion length is detrimental to cell performance. In [12], a practical 

investigation was performed by fabricating multiple mini solar cells (0.9 x 0.9 mm2) to 

examine VOC across the mc-Si wafer. A strong correlation between local VOC and minority 

carrier diffusion length was confirmed in these mini solar cells. It was concluded that a 

narrow distribution of diffusion length is desirable. Bell et al. [3-13] concluded that 

material inhomogeneities can significantly limit the voltage output, and grains with the low 

carrier lifetime determine the VOC of mc-Si cells. Warta et al. [3-14] pointed out the 

limitations of the 1D simulation tools for predicting the performance of a mc-Si solar cell 

with a defective crystal structure. A diode network model was proposed and applied to 

compare the simulated and measured mc-Si solar cell performance. Nagel et al. [3-15] used 

the diode network model developed in [3-14] to simulate the solar cell performance 

parameters on the basis of lifetime mapping on commercial cast material, neglecting the 

surface recombination and metal contact effects. In [3-16], the diode network model was 

also used to demonstrate that a 20% defective region can lead to a 30 mV reduction in VOC. 

Donolato [3-17] applied the Voronoi network model to mc-Si with columnar grains and 

showed that the grains with low carrier lifetime are responsible for the reduction in VOC. 

Isenberg et al. [3-18] showed by 2D DESSIS simulation that the network model is not 

applicable if the lateral size of defective structures is smaller than the thickness or diffusion 

length in good regions. 

In this section, an effort is made to improve the experimental and theoretical 

understanding of the impact of spatial distribution of defect inhomogeneities on fully 

processed high-performance screen-printed ribbon mc-Si solar cells. Since ribbon Si 

materials generally have an as-grown carrier lifetime below 5 μs, an attempt is made to 

enhance the gettering and passivation techniques to raise the average carrier lifetime above 

90 μs in order to increase the sensitivity and the impact of low diffusion length regions 

(bad regions). Several 4 cm2 screen-printed String Ribbon Si solar cells are fabricated and 

a simple methodology is developed to approximately determine the loss in VOC in these 

cells resulting from a small fraction of bad region mixed with a large fraction of good 

region. Model calculations are performed by developing and using simple and approximate 

analytical expressions to assess the impact of recombination intensity and area fraction of 

bad regions on VOC. Model calculations are compared with the experimental data to 

demonstrate that loss in VOC resulting from material inhomogeneity can be predicted with 

reasonable accuracy by dividing the cell into two regions (best and worst) and using this 

simple analytical model. 

30
3.2 Experiment 

In this study, 300 μm thick, p-type 3-4 Ωcm String Ribbon Si material from Evergreen 

Solar (Marlboro, MA) was used. After the initial cleaning process, the wafers were P 

diffused in a POCl3 furnace to form 40-50 Ω/sq n+ emitters. A SiNx anti-reflection coating 

with a thickness of 800 Å and a refractive index of 2.0 was deposited in a low-frequency 

plasma-enhanced chemical vapor deposition (PECVD) reactor. A commercial Al paste was 

then screen printed on the entire back surface and an Ag grid was screen printed on the 

front, followed by an anneal in a rapid thermal processing (RTP) system to (i) form an Aldoped back surface field (Al-BSF) on the rear surface, (ii) form a screen-printed Ag ohmic 

contact on the front surface, and (iii) promote the SiNx-induced hydrogen passivation of 

defects. Several 4 cm2 solar cells were fabricated on each wafer and isolated using a dicing 

saw, followed by a forming gas anneal at 400 °C for 10 min. The solar cell parameters 

were extracted by illuminated and dark current-voltage (I-V) measurements. The light 

beam-induced current (LBIC) scans were performed on selected solar cells to map the 

spatial non-uniformity of photoresponsivity using the PVSCAN5000 system [3-19] 

equipped with a 980 nm laser. Several regions on each solar cell were then selected to 

perform local light-biased internal quantum efficiency (IQE) measurements to quantify the 

difference in material quality in those regions in terms of Leff, which includes bulk and 

surface effects. 

3.3 Results and Discussion 

3.3.1 LBIC Scans and IQE Measurements 

The LBIC scan was performed on each cell to detect the material inhomogeneity and 

measure the recombination intensity of electrically active defects. Since several 4 cm2

 solar 

cells were fabricated on each large-area String Ribbon Si wafer, we selected the wafer 

containing cells with high, moderate, and low VOC to conduct our study and evaluate the 

impact of inhomogeneously distributed electrically active defects on VOC. Fig. 3-1 shows 

the LBIC maps of three String Ribbon Si solar cells on the same wafer with high, 

moderate, and low VOC. These cells were selected because they also showed maximum 

contrast between defective and defect-free regions. Table 3-2 shows the performance 

parameters of these three String Ribbon Si solar cells. The LBIC maps in Fig. 3-1 show 

that cell 1 has relatively uniform current collection over the entire cell area, compared to 

the other two, with little or no detectable bad regions. This is also reflected in the high 

area-averaged LBIC response of 0.568 A/W along with a high VOC of 616 mV and a cell 

efficiency of 15.9%. Cells 2 and 3 had a much lower VOC of 592 and 578 mV with cell 

efficiencies of 15.0% and 14.1% and corresponding area-averaged LBIC response of 0.504 

and 0.454 A/W, respectively (Table 3-2). 

31
Cell 1

VOC = 616 mV

LBIC = 0.568 A/W

Cell 2

VOC = 592 mV

LBIC = 0.504 A/W

Cell 3

VOC = 578 mV

LBIC = 0.454 A/W

0.62

0.32

Amps/Watt

A1 A3

C3

A2

C2

Fig. 3-1. LBIC maps of String Ribbon Si solar cells. High, moderate, and low VOC

cells are shown in left, center, and right sides, respectively. 

Table 3-2 Measured String Ribbon Si solar cell parameters and LBIC responses. 

Cell ID VOC 

(mV) 

JSC

(mA/cm2) FF Eff. (%) LBIC (A/W) 

1 616 33.5 0.773 15.9 0.568 

2 592 32.7 0.774 15.0 0.504 

3 578 31.4 0.778 14.1 0.454 

After the illuminated I-V measurements and LBIC scans, light-biased IQE measurements 

were taken in selected areas of these String Ribbon Si solar cells, indicated by the circles in 

Fig. 3-1. Circles drawn are for visual aid and larger than the actual spot size (~1 mm in 

diameter) used for the IQE measurements. The circled regions A1, A2, and A3 on cells 1, 

2, and 3, respectively, were selected based on the LBIC maps because they had the highest 

and nearly equal LBIC responses. These represent the best regions on the three cells. On 

the other hand, regions C2 and C3 on cells 2 and 3, respectively, gave the poorest LBIC 

response, indicating that these regions are the most defective regions on the cell. Fig. 3-2 

shows the light-biased IQE response of all five regions (A1, A2, A3, C2, and C3). As 

expected, regions C2 and C3 showed a significant degradation in the IQE response in the 

long-wavelength range (>650 nm) relative to regions A1, A2, and A3. The IQE response in 

the long-wavelength range is indicative of the combined effect of carrier lifetime and back 

surface recombination velocity (BSRV) for 300 μm thick Si cells [3-20]. Consistent with 

32
the LBIC response, there was no appreciable difference in the IQE response of regions A1, 

A2, and A3 in the long-wavelength range, supporting the fact that the best regions on the 

three solar cells (1, 2, and 3) are nearly identical, despite the significant difference in VOC

and defect inhomogeneity. 

20

30

40

50

60

70

80

90

100

550 650 750 850 950 1050 1150

Wavelength (nm)

IQE (%)

A1

.4 Extraction of Effective Diffusion Length Leff from the IQE Response 

OC is a strong function of recombination in the emitter, base, and at the surfaces. 

Ac

3

V

cording to theory, the VOC can be expressed as follows [3-21]: 

 (3-1) 

wh e k is the Boltzmann constant, T is the temperature, J er SC is the short-circuit current 

density, Job is the base components of saturation current density, and Joe is the emitter 

component of saturation current density. Since the three solar cells were selected from the 

same wafer, they should have identical emitters, and their Joe, which is a function of 

emitter doping profile and surface recombination velocity, can be assumed to be identical. 

A2

A3

C2

C3

Leff = 870 μm

Leff = 95 μm

Leff = 90 μm

Measured

Simulated

Fig. 3-2. Light-biased IQE response of selected regions (A1, A2, A3, C2, and C3) 

in String Ribbon Si solar cells and simulated (PC1D) IQE response corresponding 

to Leff values of 870, 95, and 90 μm. 

33
Thus, the Joe, which is a function of carrier lifetime and BSRV, must account for the 

majority of the difference in VOC. The Job is expressed as follows [3-22]: 

 (3-2) 

where q is the electron charge, ni is the intrinsic carrier concentration, Dn is the diffusion 

coefficient of electron, and NB is the background doping concentration. The effective 

minority carrier diffusion length, Leff, is defined as 

 (3-3) 

where S represents the BSRV, Lb (= bn D τ ) is the bulk minority carrier diffusion length, 

τb is the bulk carrier lifetime, and W is the cell thickness. According to Eq. (3-3), Leff is 

primarily a function of bulk and surface recombination for the fixed resistivity and 

thickness. 

The Leff value in any region of a solar cell can be extracted by matching the measured 

and simulated local area IQE response in the long-wavelength range (700-920 nm for 

300 μm thick cells) using the PC1D simulation program [3-23]. This was done by 

determining and/or using the realistic inputs for the PC1D simulations. For example, front 

surface recombination velocity (FSRV) was extracted from the measured IQE response in 

the short-wavelength range [3-24]. Front doping was determined by the spreading 

resistance measurement. Base contact (or series resistance) and internal conductor (or 

shunt resistance) values were obtained from illuminated I-V measurements. The internal 

diode or junction leakage current Jo2 value was extracted from the dark I-V measurement, 

assuming the second diode ideality factor of 2.0. Uniform distribution of Dn was assumed 

during the Leff extraction using PC1D. Sontag et al. [3-25] reported a factor of two lower 

Dn value in ribbon Si material compared to monocrystalline Si material. The IQE response 

in long-wavelength range represents the correct value of Leff, which is composed of Lb and 

BSRV, regardless of the Dn and τb values. Figure 3-2 also shows the match between the 

measured and simulated IQE response from all five regions, resulting in an Leff of ~870 μm 

for regions A1, A2, and A3, ~95 μm for region C2, and ~90 μm for region C3. This was 

done by fixing the BSRV at 250 cm/s and varying the Lb (or τb) in PC1D until a good 

match in the long-wavelength range was achieved. BSRV and Lb were then used to obtain 

Leff from Eq. (3-3). Note that these BSRV and Lb values do not necessarily represent the 

true BSRV and Lb values, but they still give the correct value of Leff. This is because the 

Leff is a function of both Lb (or τb) and BSRV [Eq. (3-3)]. This was validated by choosing 

different values of BSRV and performing the same procedure on IQE. As expected, 

different combinations of Lb and BSRV led to the same Leff value. 

34
Table 3-3 PC1D input parameters for String Ribbon Si solar cell simulation. 

Device Parameter Input 

Base resistivity 3.0 Ωcm 

Thickness 300 μm 

FSRV 150,000 cm/s 

Broadband reflectance 6.0% 

Front surface coating 78 nm, index = 2.0 

Rear internal reflectance 55%, diffuse 

Front doping Spreading resistance 

measurement 

Base contact 

(Series resistance) 

0.75 Ω

(0.75 Ωcm

2

) 

Internal conductor 

(Shunt resistance) 

5.0×10-5 S 

(2.0×104 Ωcm

2

) 

Jo2 2.6×10-8 A 

The Leff value of 870 μm in region A1, in conjunction with the PC1D inputs in Table 3-

3, was used to obtain the simulated solar cell I-V parameters, which agreed very well with 

the measured parameters of cell 1 (Table 3-4), which was nearly uniform. However, there 

was a significant difference between the measured solar cell parameters of the defective or 

poor solar cell 3 (Table 3-4) and the simulated cell parameters when a Leff of 870 μm was 

used for the entire cell. Table 3-4 shows that there is a difference of 37 mV in VOC. This 

difference is largely attributed to the existence of electrically active defects. 

Table 3-4 Measured and simulated solar cell parameters on String Ribbon Si using extracted 

Leff of 870 μm. 

Cell ID VOC (mV) JSC (mA/cm2) Eff. (%) 

Simulated 615 33.6 15.9 

1 616 33.5 15.9 

3 578 31.4 14.1 

The next section shows the development of a simple analytical model to calculate the 

loss in VOC from the area fraction and intensity of the recombination of the electrically 

active defects. The simulated results are then compared with the experimental data in 

Table 3-4. 

35
3.5 Theoretical and Experimental Assessment of the Impact of Electrically 

Active Defects on Solar Cell Performance 

3.5.1 Development of the Analytical Model to Assess the Loss in VOC Resulting from 

Inhomogeneity 

The non-uniform distribution of defects, such as grain boundaries, impurities, and 

dislocations, is common in mc-Si materials. Many of these defects act as localized carrier 

recombination centers resulting in spatial variation in Leff. 

The first step was to develop a simple and approximate analytical model for the 

quantitative assessment of the effect of distributed Leff on VOC. A mc-Si solar cell is 

essentially composed of a number of small solar cells (good and bad regions) in parallel, 

operating at a constant potential across the junction as a result of the electrical connection 

through grid, emitter, and substrate [3-26]. The total cell current, I, can be expressed as the 

summation of currents from local area cells: 

 (3-4) 

Assuming the two-diode model, Eq. (3-4) can be expressed as 

 (3-5) 

where V is the junction potential, Io1 and Io2 are the dark saturation current and the junction 

leakage current, respectively, and n2 is the second diode ideality factor [3-27]. Since the Io2

component generally has a much smaller contribution to total current at VOC compared to 

Io1 component, Eq. (3-5) can be expressed as 

 (3-6) 

Since Io1 is composed of Ioe and Iob, and Ioe can be independent of the defective region 

because carrier lifetime in the emitter region is generally dictated by the Auger 

recombination. Equation (6) can now be expressed as 

 (3-7) 

 (3-8) 

 (3-9) 

36
where Job,avg is the area-weighted average of Job, a is the actual area, and Ai is the area 

fraction (ai/a) of each region. Since Job is inversely proportional to the Leff [Eq. (3-2)], Eq. 

(3-9) can be written as 

 (3-10) 

The area-weighted average value of Leff (Leff,avg) can now be expressed as 

 (3-11) 

The effective Job (Job,eff) can now be determined by Eq. (3-2) using the Leff,avg. Finally, the 

Job,eff can be used as follows to obtain the VOC. 

 (3-12) 

where 

 (3-13) 

Note that Eq. (3-12) ignores the effect of Joe, which is the case for these relatively low 

VOC (<620 mV) solar cells where base quality dominates the VOC. To support this 

assumption, Joe was measured on high-resistivity (500-1000 Ωcm) n-type FZ Si using a 

photoconductance decay technique after double-sided P diffusion, SiNx deposition, and 

firing in an RTP unit without the metallization. The measured Joe of 187 fA/cm2

, without 

metallization, corresponds to a VOC of 671 mV (assuming JSC of 33.5 mA/cm2), which is 

much higher than the measured VOC of these cells. The presence of a front grid is expected 

to lower the emitter-limited value of VOC somewhat, but it will still be much higher than 

the actual VOC of these cells. 

The loss in VOC can now be calculated using the following analytical expression: 

37
 (3-14) 

where VOC (defect-free) corresponds to the VOC value in the absence of localized bad regions 

with uniformly distributed high Leff value (Leff,high=870 μm in this study). VOC(defective)

corresponds to Leff,avg and includes the combined effect of bad regions. The JSC,high and 

JSC,avg values, corresponding to Leff,high and Leff,avg were obtained from PC1D simulations 

by fixing the BSRV to the same value (250 cm/s) and adjusting the carrier lifetime to 

obtain the desired Leff for PC1D simulation. Note that the choice of BSRV does not alter 

the PC1D output or JSC significantly as long as the Leff is the same. 

3.5.2 Model Calculations to Assess the Loss in VOC of a Cell With Two Regions of 

Different Recombination Intensity 

Model calculations were performed using Eqs. (3-11), (3-13), and (3-14) to quantify the 

impact of inhomogeneous material quality on the VOC of a cell by first dividing it into two 

regions (good and bad). The Leff,high for the good region was fixed at 870 μm (Leff = 3 times 

cell thickness W), the Leff,low for bad region was varied, and the ratio of Leff,low/Leff,high was 

used as a measure of the recombination intensity in the bad region. This ratio represents 

the recombination activity in bad region, and it increases with the increase in Leff,low. The 

Leff,low values of 44, 87, 174, and 435 μm were selected for model calculation, 

corresponding to recombination intensities (Leff,low/Leff,high) of 0.05, 0.1, 0.2, and 0.5, 

respectively. To generate model curves, the area fraction of the bad region was varied from 

0 to 50% for each recombination intensity and Eqs. (3-11), (3-13), and (3-14) were used in 

sequence to calculate Leff,avg, Job,eff, and ∆VOC. JSC,high and JSC,avg were calculated from 

PC1D using Leff,high and Leff,avg as inputs. All other inputs (Table 3-3) were kept the same. 

Figure 3-3 shows the results of analytical model calculations. It shows that 10% of the 

defective area with a recombination intensity of 0.05 can reduce VOC by 28 mV of a device 

that is capable of producing VOC of 623 mV, which was obtained from Eqs. (3-12) and (3-

13) using Leff,avg of 870 μm and JSC of 33.5 mA/cm2

. Figure 3-3 also shows that we need to 

get below a 5% area fraction with a recombination intensity of 0.5 to avoid any appreciable 

loss in VOC. To determine the error associated with neglecting Joe, the VOC values were also 

calculated with Joe of 187 fA/cm2

 using the following equation: 

38
 (3-15) 

The error was found to be ±3 mV in this model calculation. 

28

17

9

2

64

58

51

42

28

35

41 46

29

26 21

16

5 7 9 11

0

10

20

30

40

50

60

70

0 10 20 30 40 50

Area of Defective Region (%)

ΔVoc (mV)

0.05

0.1

0.2

0.5

Cell 3

Cell 2

Recombination Intensity

Fig. 3-3. Calculated loss in VOC as a function of defective area with different Leff

ratio or recombination intensity. Cell was divided into two regions (good and bad). 

Base resistivity was assumed to be 3.0 Ωcm.

3.5.3 Model Calculations to Assess the Loss in VOC of a Cell With Three Regions of 

Different Recombination Intensity 

To estimate the error associated with dividing the cell into just two regions, rather than 

multiple regions, model calculations were performed to include a third region, referred to 

as the moderate region. In this calculation, the area fractions of all three regions (good, 

moderate, and bad) were varied. Each curve was obtained by fixing the area fraction of the 

bad region first and then varying the area fraction of the moderate region from 0% to 50%. 

The Leff of the good and bad regions were kept at 870 μm and 95 μm. The Leff in the 

moderate region was fixed at 400 μm and calculations were performed using Eqs. (3-11), 

(3-13), and (3-14) in sequence to obtain Leff,avg, Job,avg, and ∆VOC. The modeling results in 

Fig. 4 reveal that the bad region has the biggest influence on the loss in VOC (17 mV) even 

when its area fraction is 10%. This is because, as the moderate area fraction increases from 

39
0% to 50%, for the 10% bad area fraction, the loss in VOC increases from 17 mV to only 

24 mV. Figure 3-4 also shows that if the bad area fraction approaches or exceeds 40%, 

then it essentially dominates the loss in VOC, which reaches ~40 mV in this example. Thus, 

an approximate analysis using the best and the worst regions can give a reasonable idea of 

the loss in VOC resulting from the inhomogeneous distribution of electrically active defects. 

One can choose more than two or three regions for better accuracy using the same 

methodology. For a more exact analysis, point-by-point LBIC and IQE maps are needed 

over the entire cell area to calculate Leff,avg and Job,eff. 

41

44

21

44 42

35

39

37

40

32

30

28

35

26

24

18

29

15

20

25

30

35

40

45

0 20 40 60 80 100

Area of moderate region (%)

ΔVoc (mV)

40%

30%

20%

10%

Bad Area Fraction

Fig. 3-4. Calculated loss in VOC for each area fraction of defective region as a function 

3.5 Application of the Analytical Model to Defective Cells 

 statistical analysis of LBIC maps in Fig. 3-5 was used to assess the area fraction of 

def

of area fraction of moderate region. Base resistivity was assumed to be 3.0 Ωcm.

.4

A

ective regions. The LBIC output gives a map or range of spectral responses in A/W, 

which was divided into 16 equal size bins. Figure 3-5 shows the resulting histograms of the 

LBIC response for the three cells. Next, each cell was divided into two regions: good 

region with a spectral response ranging from 0.48 to 0.62 A/W and bad region with a 

spectral response ranging from 0.32 to 0.46 A/W. The area fraction of the two regions was 

determined by 

40
 (3-16)

where Bi is the count in the i bin, N is the number of bins in that region (e.g., 0.32-0.46 

A/W for region 2), and K is the total number of bins (=16 in this study). This analysis gave 

the area fractions of 0.02, 0.19, and 0.38 associated with the bad regions for cells 1, 2, and 

3, respectively. Note that the grid coverage (6% of cell area) was removed from the 

histogram in Fig. 5 since there is no carrier generation below the grid. It is important to 

note that the same procedure can be used to divide the LBIC output into multiple regions 

instead of just two. 

Recall that the light-biased IQE response in Fig. 3-2 showed a significant degradation 

in the long-wavelength IQE response of regions C2 and C3, relative to regions A2 and A3, 

and the measured and simulated IQE match gave Leff values of 95 μm and 90 μm for 

regions C2 and C3, respectively, as opposed to 870 μm for regions A2 and A3. According 

to Eq. (3-2), this results in significantly higher Job values of 10-12 to 10-13 A/cm2 in 

regions C2 and C3 relative to the good regions A2 and A3 where J of 10-14 A/cm2

 is 

obtained. 

Using Eq. (3-11) and a Leff,high of 870 μm and Leff,low of 95 μm with an area fraction of 

81% and 19% for the good and bad regions for cell 2 gives a Leff,avg value of 342 μm. 

Using the Leff values of 870 and 342 μm along with the PC1D inputs in Table 3-3, JSC

values of 33.6 mA/cm2

 (JSC,high) and 32.6 mA/cm2 (JSC,avg) are obtained from the PC1D 

simulations. Note that these JSC values are in good agreement with the measured values of 

0

50000

100000

150000

200000

<0.32 0.420.400.380.360.34 0.44 0.620.600.580.560.540.520.500.480.46

LBIC Response (A/W)

Counts

Cell 1

Cell 2

Cell 3

Fig. 3-5. Histograms of LBIC response shown in Fig. 1. The Ag grid 

coverage effect (6.0% of cell area) was removed from <0.32 A/W bin. 

Maximum LBIC response was 0.62 A/W for all three cells. 

41
33.5 mA/cm2

 and 32.7 mA/cm2 for cell 1 and 2, respectively (Table II). Finally, using 

these JSC and Leff values in Eq. (14) gives a ∆VOC of 25 mV, which is in good agreement 

with the measured VOC difference of 24 mV between cell 1 and 2 (Table 3-2). 

Cell 3 was also analyzed using the same methodology described above. An Leff,high of 

870 μm and Leff,low of 90 μm with a good and bad area fraction of 62% and 38% for cell 3 

gave a Leff,avg value of 203 μm, which resulted in a JSC,avg value of 31.6 mA/cm2 from the 

PC1D simulation. Note that this JSC value is in good agreement with the measured value of 

31.4 mA/cm2 for cell 3 (Table 3-2). Equation (3-14) gave a ∆VOC of 39 mV, which is again 

in good agreement with the measured VOC difference of 38 mV between cells 1 and 3 

(Table II). Cells 2 and 3 are mapped in Fig. 3-3 to show the agreement between cell data 

and model calculations. Even though the model and the procedure outlined above are valid 

for multiple-region analysis, a simple two-region analytical model gave a fairly good idea 

of the loss in VOC resulting from the material inhomogeneity in these high-performance 

screen-printed mc-Si solar cells. 

3.6 Conclusions 

The effect of electrically active defects on VOC was quantified and evaluated through a 

simple and approximate analytical model. Model calculations were performed to reveal the 

relationship between the area fraction, recombination intensity of defective region, and the 

loss in VOC. In this study, three String Ribbon Si solar cells were used to test this model 

and understand the impact of defect inhomogeneity. A two-region analysis using the model 

developed in this study gave a reasonable estimate of the loss in VOC resulting from defect 

inhomogeneity. LBIC and IQE measurements were used to estimate the area fraction and 

recombination intensity of the defective regions. Model calculations showed that a cell 

with a 38% area fraction of bad region with a recombination intensity of 0.1 can reduce the 

VOC by 39 mV. This agreed well with the experimental data for cell 3, which showed a loss 

in VOC of 38 mV. There was also a good agreement between the experimental data and 

model calculations, which showed that a 19% area fraction of bad region with a 

recombination intensity of 0.11 can reduce the VOC by 25 mV. Model calculations revealed 

that in order to keep the loss in VOC below 5 mV in high-performance devices, the area 

fraction of the defective region should be below 20% with a recombination intensity of 0.5 

or greater. 

42
3.7 References 

[3-1] United States Department of Energy, “Solar Photovoltaic,” Feb. 2005, 

http://www.eia.doe.gov/cneaf/solar.renewables/page/solarphotv/solarpv.html

[3-2] R.O. Bell and J.P. Kalejs, “Growth of silicon sheets for photovoltaic applications,” J. 

Mater. Res., vol. 13, pp. 2732-2739, Oct. 1998. 

[3-3] J.I. Hanoka, “An overview of silicon ribbon growth technology,” Sol. Energy Mater. 

Sol. Cells, vol. 65, pp. 231-237, Jan. 2001. 

[3-4] J. Jeong, A. Rohatgi, M.D. Rosenblum and J.P. Kalejs, in Proc. 28th IEEE 

Photovoltaic Specialists Conference, 2000, pp. 83-86. 

[3-5] J. Jeong, Y.H. Cho, A. Rohatgi, M.D. Rosenblum, B.R. Bathey and J.P. Kalejs, in 

Proc. 29th IEEE Photovoltaic Specialists Conference, 2002, pp. 250-253. 

[3-6] F. Duerinckx and J. Szlufcik, “Defect passivation of industrial multicrystalline solar 

cells based on PECVD silicon nitride,” Sol. Energy Mater. Sol. Cells, vol. 72, pp. 231-

246, Apr. 2002. 

[3-7] A. Rohatgi, V. Yelundur, J. Jeong, A. Ebong, M.D. Rosenblum, and J.I. Hanoka, 

“Fundamental understanding and implementation of Al-enhanced PECVD SiNx 

hydrogenation in silicon ribbons,” Sol. Energy Mater. Sol. Cells, vol. 72, pp. 117-126, 

Oct. 2002. 

[3-8] V. Yelundur, A. Rohatgi, A. Ebong, A.M. Gabor, J.I. Hanoka, and R.L.Wallace, “Alenhanced PECVD SiNx induced hydrogen passivation in string ribbon silicon,” J. 

Electron. Mater., vol. 30, pp. 526-531, May 2001. 

[3-9] J. Jeong, M.D. Rosenblum, J.P. Kalejs, and A. Rohatgi, “Hydrogenation of defects in 

edge-defined film-fed grown aluminum-enhanced plasma enhanced chemical vapor 

deposited silicon nitride multicrystalline silicon,” J. Appl. Phys., vol. 87, pp. 7551-

7557, May 2000. 

[3-10] B.L. Sopori, “Inhomogeneities in silicon solar cells and their influence on cell 

performance - an experimental study,” in Proc. 14th IEEE Photovoltaic Specialists 

Conference, 1980, pp. 606-611. 

[3-11] P.E. Mijnarends, G.J.M. Janssen, and W.C. Sinke, “The effect of material 

inhomogeneities on the characteristics of semicrystalline silicon solar cells: the second 

diode,” Sol. Energy Mater. Sol. Cells, vol. 33, pp. 345-360, July 1994. 

[3-12] R. Baldner, H. Lautenschlager, C. Schetter, R. Schindler, and W. Warta, “Open 

circuit voltage losses in multicrystalline silicon: an investigation by mini solar cells,” in 

Proc. 25th IEEE Photovoltaic Specialists Conference, 1996, pp. 641-644. 

[3-13] R.O. Bell, M. Prince, F.V. Wald, W. Schmidt, and K.D. Rasch, “A comparison of 

the behavior of solar silicon material in different production processes,” Sol. Energy 

Mater. Sol. Cells, vol. 41/42, pp. 71-86, June 1996. 

[3-14] W. Warta, J. Sutter, B.F. Wagner, and R. Schindler, “Impact of diffusion length 

distributions on the performance of mc-silicon solar cells,” in Proc. 2nd World 

Conference on Photovoltaic Energy Conversion, 1998, pp. 1650-1653. 

43
[3-15] H. Nagel, A.G. Aberle, and S. Narayanan, “Method for the evaluation of the 

influence of gettering and bulk passivation on non-uniform block-cast multicrystalline 

Si solar cells,” Di_us. Defect Data B, Solid State Phenom., vol. 67-68, pp. 503-508, 

1999. 

[3-16] B.L. Sopori, W. Chen, J. Gee, and K. Jones, “On the performance limiting behavior 

of defect clusters in commercial silicon solar cells,” in Proc. 2nd World Conference on 

Photovoltaic Energy Conversion, 1998, pp. 152-155. 

[3-17] C. Donolato, “Voronoi network modelling of multicrystalline silicon solar cells,” 

Semicond. Sci. Technol., vol. 15, pp. 15-23, Jan. 2000. 

[3-18] J. Isenberg, J. Dicker, and W. Warta, “Analysis of the effect of diffusion length 

distributions on global solar cell parameters by simplified 2D modeling,” in Proc. 17th

European Photovoltaic Solar Energy Conference and Exhibition, 2001, pp. 1571-1574. 

[3-19] K.F. Carr, N. Carison, P. Weitzmann, B.L. Sopori, C. Marshall, and L. Allen, 

“Characterization of silicon solar cells and substrates with the PVScan 5000,” in AIP 

Conference Proc. of 13th NREL Photovoltaic Program Rev., 1996, pp. 553-557. 

[3-20] V. Meemongkolkiat, M. Hilali, and A. Rohatgi, “Investigation of RTP and belt fired 

screen printed Al-BSF on textured and planar back surfaces of silicon solar cells,” in 

Proc. 3rd World Conference on Photovoltaic Energy Conversion, 2003, pp. 1467-1470. 

[3-21] M.A. Green, Solar Cells, Kensington: UNSW, 1986. 

[3-22] P.A. Basore, “Extended spectral analysis of internal quantum efficiency,” in Proc. 

23rd IEEE Photovoltaic Specialists Conference, 1993, pp. 147-152. 

[3-23] V. Meemongkolkiat, M. Hilali, K. Nakayashiki, and A. Rohatgi, “Process and 

material dependence of Al-BSF in crystalline Si solar cells,” in Technical Digests of 

14th International Photovoltaic Science and Engineering Conference, 2004, pp. 401-

402. 

[3-24] M.M. Hilali, A. Rohatgi, and S. Asher, “Development of screen-printed silicon solar 

cells with high fill factors on 100 /sq emitters,” IEEE Trans. Electron Devices, vol. 51, 

pp. 948-955, June 2004. 

[3-25] D. Sontag, G. Hahn, P. Geiger, P. Fath, and E. Bucher, “Two-dimensional 

resolution of minority carrier diffusion constants in different silicon materials,” Sol. 

Energy Mater. Sol. Cells, vol. 72, pp. 533-539, 2002. 

[3-26] R. Sinton, “Predicting multicrystalline solar cell efficiency from lifetime measured 

during cell fabrication,” in Proc. 3rd World Conference on Photovoltaic Energy 

Conversion, 2003, pp. 1028-1031. 

[3-27] S.M. Sze, Semiconductor Devices Physics and Technology, New York: John Wiley 

& Sons, 1985. 

44
4. Resistivity and Lifetime Variation along Commercially Grown 

Ga- and B-Doped Czochralski Si Ingots and Its Effect on LightInduced Degradation and Performance of Solar Cells 

4.1 Introduction 

It is well known that solar cells fabricated on conventional B-doped Czochralski (Cz) Si 

suffer from degradation caused by the illumination or injection of carriers [4-1,4-2]. It has 

also been established that this light-induced degradation (LID) effect results from the 

presence of B and O simultaneously in Si [4-3,4-4]. Therefore, the LID effect can be 

removed by eliminating either B or O from Si. Several alternatives have been suggested 

and attempted in the literature to avoid LID in Cz Si including the 1) use of alternative 

dopants such as P for n-type and Ga for p-type Cz Si [4-3,4-5,-6] 2) reduction of interstitial 

oxygen to an acceptable level by growing Magnetic Cz [5] Si 3) use of higher resistivity Bdoped Cz [4-2,4-7] and 4) process optimization [4-4,4-8 - 4-13]. Doping the Si ingot with 

Ga has advantages over other methods as it provides complete elimination of LID without 

modifying the cell structure or processing equipment. However, there are some drawbacks 

associated with the implementation of Ga. First, use of Ga dopant gives rise to a 

complication of managing the silicon feedstock [4-14]. Most Cz Si manufacturers for solar 

cells application rely on externally supplied remelt or potscrap. As Ga-doped remelt is not 

available on the open market, this situation likely necessitates the recycling of internally 

produced remelt, subsidized with virgin Si. Relying on virgin suppliers is not a costeffective option and creates a significant risk since the availability and, more importantly, 

the cost of virgin Si fluctuates greatly. The other basic drawback of using Ga as a dopant is 

the low segregation coefficient of Ga in Si (k=0.008). This results in a much wider 

variation in resistivity along the Ga-doped ingot. If the solar cell fabrication process cannot 

tolerate this wider resistivity distribution, a significant yield loss in the crystal growth 

process would result in a substantial cost addition. Many research groups have tried to 

investigate the solar cell performance as a function of resistivity or the position in the ingot 

[4-11,4-15,4-16]. However, such relationship has never been established for the widely 

manufactured screen-printed solar cells on Si wafers obtained from B- and Ga-doped 

commercial Cz ingots grown in the same puller. Therefore, in this study, B- and Ga-doped 

Cz ingots were grown in an industrial environment at Shell solar industries, and the Cz Si 

wafers from different locations along the Ga- as well as the B-doped ingots were analyzed. 

Both the ingots were targeted to have resistivity of ~1 Ω-cm. In addition to the wafers from 

these two ingots, some B-doped thin wafers with higher resistivity (~4.3 Ω-cm) were 

included in the study to explore the reduction in the LID effect as proposed in the literature 

[4-17,4-18]. All three ingots were grown using the exact same growth method and 

equipment. The bulk lifetime in all the samples was determined by the contactless 

photoconductance measurement [4-19,4-20]. Manufacturable screen-printed solar cells 

were fabricated and analyzed using light IV-measurement. 

45
4.2 Experimental Procedures 

The Si wafers used in this study were taken from different locations along the two Cz 

Si ingots: six locations from the 950 mm long, ~1 Ω-cm, B-doped ingot and nine locations 

from the 925 mm long, 0.5-2.5 Ω-cm, Ga-doped ingot were selected (Table 4-1). 

Additional wafers were taken from a higher resistivity B-doped ingot (~4.3 Ω-cm) and 

thinned down to ~230 μm from ~290 μm. Table I summarizes all the wafers used in the 

study. The lifetime was measured on each sample both in the as-grown state and after the 

phosphorus emitter diffusion. The post-diffusion lifetime was measured on each sample 

after POCl3 diffusion at ~880°C followed by etching of the sample down to Si bulk. The 

lifetime measurements were performed after 1) 200°C anneal to remove any LID effect [4-

1] and 2) light-soaking for >20 hrs to obtain the stabilized lifetime after LID. The surface 

was passivated by iodine/methanol solution during the lifetime measurements. 

 Screen-printed Al-back surface field (BSF) solar cells (4 cm2

) were fabricated on all 

the wafers in Table 4-1 using an industrial process. First, the samples were textured in an 

alkaline etch and then POCl3 diffused to obtain a ~45 Ω/sq emitter. Subsequently, SiNx

AR-coating was deposited on the front. All the samples were then subjected to full-area Al 

screen-printing on the backside, followed by Ag gridline printing on the front. The samples 

were then co-fired using rapid thermal processing. No special heat treatment was 

performed to minimize LID as proposed in the literature [4-4,4-8 – 4-13]. 

The I-V measurements were taken after annealing the cells at 200°C to remove the LID 

effect and to determine the cell performance without the LID. The I-V measurements were 

repeated on all the cells after light soaking them for >20 hrs to obtain the stabilized cell 

performance after LID. 

4.3 Results and Discussion 

4.3.1 Growth of Ga-doped Cz Ingots 

The crystal growth process appears to be transparent to the use of Ga dopant because 

same method and equipment was used successfully for B- and Ga-doped Cz growth. In 

comparison with B-doped Cz Si, the growth of Ga-doped Cz Si does not require an extra 

effort on the growth control, melt contamination or maintaining structure loss. Maintaining 

the rate of structure loss is especially important for cost control since it is typically the 

highest loss category in a Cz crystal growth process. Consequently, no changes to the 

crystal growth process were necessary to achieve similar yields on Ga-doped ingot. 

4.3.2 Resistivity Distribution Along Commercially Grown B-doped and Ga-doped Cz 

Ingots 

Table 4-1 shows that the growth of low-resistivity B-doped ingot provides samples with 

a tight resistivity control, ranging from 0.87 Ω-cm to 1.22 Ω-cm. However, the resistivity 

variation is much larger (0.57-2.54 Ω-cm) in the case of the Ga-doped ingot compared to 

the B-doped ingot. The resistivity decreases appreciably from seed to tail end of the Ga46
doped ingot because of the low segregation coefficient of Ga (k=0.008) compared to B 

(k=0.8) in Si. 

Table 4-1. Description of Cz Si samples used in the study. 

Ingot Thickness Tail end Seed end 

Low-ρ Location 1 2 3 4 5 6 

Bdoped 

290μm 

ρ (Ωּcm) 0.87 0.82 0.90 0.95 1.00 1.22 

Ga- Location 1 2 3 4 5 6 7 8 9 

doped 290μm ρ (Ωּcm) 0.57 0.63 0.84 0.99 1.19 1.46 1.82 2.17 2.54 

Hi-ρ

Bdoped 230μm ρ (Ωּcm) 4.3 

4.3.3 As-grown and post-diffusion lifetime in wafers from B- and Ga-doped Cz ingots 

4.3.3.1 B-doped ingots 

The as-grown and post-diffusion lifetimes for the B-doped Cz samples are summarized in 

Figure 1a and 1b, respectively. The low-resistivity B-doped ingot showed a very tight 

distribution of the lifetime except at the seed end, where the lifetime was somewhat lower 

probably due to the swirl defects that occur in macroscopically dislocation-free Si with a 

high density of point defects [4-21]. All the bulk lifetime measurements in Figure 4-1 were 

performed at an injection level of 2x1014 cm-3. 

1

10

100

1000

1 2 3 4 5 6 4.3 Ω-cm

low-ρ B-doped ingot High-ρ

B-doped

ingot

Lifetime (μs)

Annealed

Light-soaked

Tail end

0.87 Ω-cm Seed end

1.22 Ω-cm

1

10

100

1000

1 2 3 4 5 6 4.3 Ω-cm

low-ρ B-doped ingot High-ρ

B-doped

ingot

Lifetime (μs)

Annealed

Light-soaked

Tail end

0.87 Ω-cm Seed end

1.22 Ω-cm

 a) b) 

Figure 4-1. (a) As-grown and (b) post-diffusion lifetime before and after LID on wafers 

from different locations in the low- and high-resistivity B-doped Cz ingots. 

47
Figure 4-1 shows that the phosphorus diffusion enhances the light-soaked lifetime 

significantly. This is attributed to 1) impurity gettering by phosphorus diffusion [4-22] and 

2) the reduction of metastable defects (responsible for LID) by high-temperature treatment. 

For this reason, the lifetime in the finished cell correlates better with the post-diffusion 

lifetime rather than the as-grown lifetime. To consider the effectiveness of the high 

temperature process in the reduction of LID, it is useful to first calculate the normalized 

metastable defect concentration, defined as [4-23]. This can be calculated 

from the lifetime in annealed and light-soaked states as follows: 

N = υσ Ntthnt

*

light soaked annealed

Nt

τ τ

1 1 = −

−

* (4-1) 

Nt

*

 was calculated using eq. (4-1) and the measured lifetime before and after LID at an 

injection level of 2x1014 cm-3 for all samples. The Nt

*

 values for as-grown and postdiffusion states are shown in Figure 4-2. Smaller value of Nt

*

 reflects lower concentration 

of metastable defects (Nt) that are responsible for LID. Our data shows that the phosphorus 

diffusion reduced the metastable defect concentration effectively by a factor of 2.8-3.1 in 

the low resistivity B-doped samples and by a factor of 2.4 in the high-resistivity B-doped 

sample. In spite of this respectable improvement, the LID managed to lower the post 

diffusion lifetime by a factor of 15-20 (from ~300-400 μs to ~20 μs) in the low-resistivity 

(~1 Ω-cm) B-doped wafers (Figure 1b). The high-resistivity B-doped sample, on the other 

hand, exhibited a much weaker LID effect due to lower B concentration, resulting in about 

a factor of three degradation in lifetime from ~500 μs to ~170 μs. 

0.001

0.01

0.1

1

1 2 3 4 5 6 4.3 Ω-cm

low-ρ B-doped ingot High-ρ

ingot

t N * (1/μs) t

As-grown

Post diffusion

Tail end

0.87 Ω-cm Seed end

1.22 Ω-cm B-doped

Figure 4-2. Normalized metastable defect concentration before and after diffusion process in 

B-doped Cz wafers. 

48
4.3.3.2 Ga-doped ingot 

The as-grown and post-diffusion lifetimes measured at the injection level of 2x1014 cm3

 for the Ga-doped Cz samples are summarized in Figure 4-3a and 4-3b, respectively. 

Unlike the B-doped wafers, Ga-doped wafers did not show any LID. However, the postdiffusion lifetime in Ga-doped ingot varied significantly from ~100 to ~1000 μs from seed 

to tail end partly because of the variation in the resistivity. Both resistivity and lifetime 

increased gradually from tail to seed end except for a drop at the seed end for the same 

reason (swirl defects) as the B-doped ingot. Notice that the phosphorus diffusion improved 

the lifetime of Ga-doped wafers by about a factor of 1.3 due to gettering of impurities. 

1

10

100

1000

123456789

Ga-doped ingot

Lifetime (μs)

Annealed

Light-soaked

Tail end

0.57 Ω-cm 

Seed end

2.54 Ω-cm 

1

10

100

1000

123456789

Ga-doped ingot

Lifetime (μs)

Annealed

Light-soaked

Tail end

0.57 Ω-cm 

Seed end

2.54 Ω-cm 

a) b) 

Figure 4-3. (a) As-grown and (b) post-diffusion lifetime before and after LID on wafers 

from different locations in Ga-doped Cz ingot. 

4.4 LID and Performance of Screen-Printed Solar Cells 

4.4.1 B-doped Ingots 

Simple manufacturable 4 cm2

 screen-printed cells were fabricated with texturing, AlBSF and SiNx PECVD single-layer AR coating (Section 2). The efficiency of the solar 

cells fabricated on wafers taken from different locations on B-doped ingots is plotted in 

Figure 4-4. Both annealed (no LID) and light-soaked (after LID) states are included in the 

figure to assess the LID effect. 

The efficiency prior to the LID in the low-resistivity B-doped ingot was quite uniform 

(~16.7%), except at the seed end where the efficiency dropped slightly to 16.5%. This is 

entirely consistent with the lifetime data in Figure 4-1b, which showed fairly uniform 

lifetime except at the seed end. However, the efficiency of all the low-resistivity B-doped 

cells decreased significantly by about 1.1% absolute after the light soaking, resulting in a 

final efficiency of only ~15.6%. This is also consistent with Figure 4-1b which shows that 

after the diffusion and light-soaking, the lifetime in all the wafers dropped to ~ 20 μs after 

LID. The high-resistivity (4.3 Ω-cm) thin (230 μm) B-doped Cz cell gave an efficiency of 

49
14.5

15.0

15.5

16.0

16.5

17.0

17.5

1 2 3 4 5 6 4.3 Ω-cm

 17%, which is ~0.3% better than the low-resistivity B-doped cells. In addition, the LID 

effect was substantially reduced with an efficiency loss of ≤0.6% absolute as opposed to 

1.1% for the low-resistivity thick cells, resulting in the stabilized efficiency of 16.4%. The 

reduced LID effect is attributed to 1) the reduction of B concentration in the higher 

resistivity material, resulting in fewer LID traps and 2) higher diffusion length to thickness 

ratio due to thinner material, resulting in decreased sensitivity to loss in diffusion length. 

Thus, higher resistivity and thinner material provides another strategy for reducing LID in 

B-doped Cz cells. 

Device simulations were performed using PC1D program [4-24] to establish that the 

impact of LID on the cell efficiency is entirely based on lifetime degradation at maximum 

power point (MPP). Some of the key inputs used for PC1D simulation are summarized in 

Table 4-2. The simulated efficiency as a function of lifetime for a 300 μm thick 1.0 Ω-cm 

substrate and 230 μm thick 4.3 Ω-cm are plotted in Figure 4-5. Due to a strong asymmetry 

of the capture time constants for electrons and holes (τn/τp~0.1), specific for boron-oxygenrelated recombination center [4-13], the bulk lifetime of a p-type material strongly depends 

on the injection level. This dependence of lifetime on injection level needs to be taken into 

account to assess the accurate and full impact of LID on cell efficiency. Therefore, first we 

measured the lifetime as a function of injection level in these samples as shown in Figure 

4-6. This was done on diffused samples after the emitter was removed. Figure 4-6 confirms 

the asymmetric capture time constant because lifetime does vary strongly with injection 

level after LID. Notice that prior to LID, lifetime is high and is not a strong function of 

injection level. This measured lifetime data was used in conjunction with PC1D calculation 

of efficiency as a function of lifetime (Figure 4-5) to assess the loss in efficiency due to 

LID. To do this accurately, the loss in bulk lifetime was determined at the MPP. The 

device simulations were performed to determine the approximate injection level at the 

low-ρ B-doped ingot High-ρ

B-doped

ingot 

Annealed

Light-soaked

Tail end

0.87 Ω-cm 

Seed end

1.22 Ω-cm

Efficiency (%)

Figure 4-4. Screen-printed Al-BSF solar cells efficiency of samples from different 

locations from low- and high-resistivity B-doped Cz ingots. 

50
injection level is not completely uniform along the device thickness. Carrier concentration 

profiles reveal that in the 1 Ω-cm cell, the injection level varied from 3x1012 cm-3 to 

1.2x1013 cm-3 at MPP for 10-30 μs bulk lifetime. Figure 6 shows that in spite of this large 

variation in injection level, the lifetime should remain nearly constant at about 15 μs. 

Similarly for the 4.3 Ω-cm cell, the injection level at MPP was found to vary from 2.0x1013 

to 3.1x1013 cm-3 for 80-200 μs bulk lifetime, resulting in a nearly constant lifetime of 105 

μs at the MPP. These lifetime values were used in Figure 4-5 to obtain the predicted cell 

efficiency after LID of 15.7% for 1.0 Ω-cm and 16.6% for 4.3 Ω-cm cells. These 

efficiencies were in good agreement with the experimental data in Figure 4 (~15.6% for 

1.0 Ω-cm and 16.4% for 4.3 Ω-cm Cz). Note that in Figure 4-5, if the measured lifetime at 

an injection level of 2x1014 cm-3 was used (which is closer to the open circuit voltage (Voc) 

condition of the cells), the simulation would have underestimated the LID effect on 

efficiency due to the strong lifetime dependence on injection level. Figure 4-5 also shows 

that before LID, predicted efficiencies are ~16.7% and ~17.0% for 1.0 Ω-cm and 4.3 Ω-cm 

cells, which is again in good agreement with the experimental data even though the 

measured lifetime at 2x1014 cm-3 was used. This is because, prior to LID, measured 

lifetime is not a strong function of injection level. Thus, in order to predict the LID effect 

by simulation, one needs to recognize the strong injection level dependence of lifetime and 

use the measured lifetime at the injection level around the peak power point. Lifetimes in 

solar cell materials are often measured at the Voc condition or even at higher injection level 

(≥1x1015 cm-3) to avoid trapping effects on lifetime data [4-25]. 

15

15.5

16

16.5

17

0 100 200 300 400 500 600

Lifetime (μs)

Efficiency (%)

1 Ω-cm (300 μm thick)

4.3 Ω-cm (230 μm thick)

0.35%

0.99%

(3) After LID at MPP

(2) After LID at injection level of 2x1014 cm-3 (1) Before LID

Figure 4-5. Simulated solar cell efficiency as a function of lifetime for a) 1.0 Ω-cm 300 μm 

thick Si substrate and b) 4.3 Ω-cm 230 μm thick Si substrate. Simulated curves are used to 

show the predicted cells efficiencies based on measured lifetime (1) before LID, (2) after 

LID but at injection level of 2x1014 cm-3, and (3) after LID at an injection level at MPP. 

51
10

100

1000

1.E+13 1.E+14 1.E+15 1.E+16

Injection level (cm-3)

Lifetime (μs)

1 Ω-cm, annealed 1 Ω-cm, light-soaked 4.3 Ω-cm, annealed 4.3 Ω-cm, light-soaked Guideline

Figure 4-6. Measured bulk lifetime as a function of injection level, before and after LID, for 

~1 Ω-cm and ~4.3 Ω-cm B-doped Cz. 

Table 4-2. PC1D input parameters for screen-printed cell simulation. 

Device Parameter Input 

Front surface Textured, 54.74°, 3.535 μm 

Front surface reflectance From measurement 

Broadband reflectance 6.75-7.00 % 

Rear internal reflectance 60%, diffuse 

Base contact 0.75-0.80 Ω

Internal conductor 1×10-4 S 

Jo2 3.2-3.5×10-8 A 

Front doping 45 Ω/sq, Erfc 

FSRV 95,000-120,000 cm/s 

BSRV 750 cm/s for 1.0 Ω-cm 

150 cm/s for 4.3 Ω-cm 

52
4.4.2 Ga-doped Ingot 

The variation in efficiency of the screen-printed Al-BSF solar cells fabricated on wafers 

from different ingot locations of Ga-doped ingot is plotted in Figure 4-7. Both annealed 

(no LID) and light-soaked (after LID) states are included in the figure. 

14.5

15.0

15.5

16.0

16.5

17.0

17.5

123456789

Ga-doped ingot

Efficiency (%)

Annealed

Light-soaked

Tail end

0.57 Ω-cm 

Seed end

2.54 Ω-cm 

Figure 4-7. Screen-printed Al-BSF solar cells efficiency of samples from different locations 

from Ga-doped Cz ingot. 

Unlike the B-doped ingots, the efficiency spread in Ga-doped ingot prior to the LID is 

somewhat larger (16.8%-17.3%), with the higher resistivity seed end producing slightly 

higher efficiency. This variation in efficiency is generally acceptable for production and is 

within the range of process-induced effects. Detailed analysis of cell parameters in Figure 

8 shows that, despite the very wide variation in resistivity over the entire length of the 

ingot, the spread in Ga-doped cell efficiency is reduced because of the increase in Voc and 

the decrease in short circuit current density (Jsc) as the resistivity decreases. Note that 

spread in current due to resistivity variation in Ga-doped cells may add to the spread in 

current due to the processing effects. However, if the cells are binned according to current 

and efficiency combination, rather than just current, then it will be possible to separate 

different resistivity cells in appropriate bins, to avoid cells’ voltage mismatch losses, prior 

to module fabrication. Moreover, there is essentially no LID observed in the Ga-doped 

cells, resulting in ~1.5% higher absolute efficiency after light soaking relative to the lowresistivity B-doped Cz ingot. This gap in stabilized efficiency was reduced to ~0.7% when 

higher resistivity and thin B-doped Cz was used. These results show that the Ga-doped Cz 

ingot offers great potential for higher stabilized Cz cell performance (≥ 17%). In addition, 

a high-quality Ga-doped ingot can be grown in the same puller used for the B-doped ingot, 

without any modification. 

53
618

620

622

624

626

628

630

123456789

Ga-doped ingot

Voc (mV)

33.5

34

34.5

35

35.5

36

36.5

Jsc (mA/cm2

)

Voc Jsc

Tail end

0.57 Ω-cm 

Seed end

2.54 Ω-cm 

Figure 4-8. Jsc and Voc as a function of ingot position in Ga-doped ingot. 

Table 4-3 summarizes the average efficiency of solar cells from the three Cz Si ingots. 

These data demonstrate the potential of using Ga-doped Cz instead of B-doped Cz. 

Table 4-3. Summary of averaged efficiency from different Cz ingots. 

Dopant Efficiency (%) 

Type 

Resistivity 

(Ω-cm) Annealed Stabilized 

~ 1.0 16.7 15.6 

Boron 

~4.3 17.0 16.4 

Gallium 0.57-2.54 17.1 17.1 

4.5 Conclusions 

This paper demonstrates the potential of using Ga dopant instead of B in p-type Cz Si 

to achieve high efficiency manufacturable screen-printed cells with no LID. Despite a large 

resistivity variation (0.57-2.54 Ω-cm) in the Ga-doped Cz ingot resulting from a small 

segregation coefficient of Ga in Si, the absolute efficiency of screen-printed Al-BSF solar 

cells was found to vary by <0.5% absolute over the entire length of the 925 mm long ingot. 

This is the result of the competing effect of increasing Voc and decreasing Jsc as the 

resistivity decreases. In the 1 Ω-cm B-doped Cz, lifetimes decreased from 300-400 to ~20 

μs after LID. In the 4.3 Ω-cm B-doped Cz, lifetimes decreased from 500 μs to 170 μs after 

LID. In the Ga-doped ingot, lifetimes were in the range of 100-1000 μs from seed to tail 

end, respectively, and showed no LID at all. This resulted in ~1.5% higher average 

54
stabilized efficiency compared to the cells made on a 1 Ω-cm B-doped Cz and about 0.7% 

higher than efficiency with respect to 4.3 Ω-cm thin B-doped cells. These results were 

found to be in good agreement with device simulations performed using the measured 

lifetime at the injection level at MPP. 

The use of thinner and high-resistivity B-doped Cz lessened the detrimental effect of 

LID. However, the LID remained appreciable and accounted for ~0.6 % absolute reduction 

in efficiency. Ga doping completely eliminated the LID in Cz cells and gave ≥ 17% 

efficient screen-printed solar cells. 

4.6 References 

[4-1] Fischer H, Pschunder W. Investigation of photon and thermal induced changes in 

silicon solar cells. In Conference Record, 10th IEEE Photovoltaic Specialists 

Conference. Palo Alto, November 1973; 404-411. 

[4-2] Knobloch J, Glunz SW, Henninger V, Warta W, Wettling W. 21% efficient solar 

cells processed from Czochralski grown silicon. In Conference Proceedings, 13th

European Photovoltaic Solar Energy Conference. Nice, October 1995; 9-12. 

[4-3] Schmidt J, Aberle AG, Hezel R. Investigation of carrier lifetime instabilities in 

Cz-grown silicon. In Conference Record, 26th IEEE Photovoltaic Specialists 

Conference. Anaheim September/October 1997; 13-18. 

[4-4] Glunz SW, Rein S, Warta W, Knobloch J, Wettling W. On the degradation of Czsilicon solar cells. In Conference Proceedings, 2nd World Conference on 

Photovoltaic Solar Energy Conversion. Vienna, July 1998; 1343-1346. 

[4-5] Glunz SW, Rein S, Knobloch J, Wettling W, Abe T. Comparison of boron- and 

gallium-doped p-type Czochralski silicon for photovoltaic application. Progress in 

Photovoltaics 1999; 7: 463-469. 

[4-6] Saitoh T et al. Light degradation and control of low-resistivity Cz-Si solar cells –

An international joint research-. In Technical Digest, 11th International 

Photovoltaic Science and Engineering Conference. Sapporo, September 1999; 

553-556. 

[4-7] Knobloch J, Glunz SW, Biro D, Warta W, Schäffer E, Wettling W. Solar cells 

with efficiencies above 21% processed from Czochralski grown silicon. In 

Conference Record, 25th IEEE Photovoltaic Specialists Conference. Washington, 

May 1996; 405-408. 

[4-8] Nagel H, Merkle A, Metz A, Hezel R. Permanent reduction of excess-carrierinduced recombination centers in solar grade Czochralski silicon by a short yet 

effective anneal. In Conference Proceedings, 16th European Photovoltaic Solar 

Energy Conference. Glasgow, May 2000; 1197-1200. 

[4-9] Rein S, Knobloch J, Glunz SW. Analysis of the high-temperature improvement of 

Cz-silicon. In Conference Proceedings, 16th European Photovoltaic Solar Energy 

Conference. Glasgow, May 2000; 1201-1205. 

[4-10] Schmidt J, Cuevas A. Progress in understanding and reducing the light 

degradation of Cz silicon solar cells. In Conference Proceedings, 16th European 

Photovoltaic Solar Energy Conference. Glasgow, May 2000; 1201-1205. 

55
[4-11] [-Glunz SW, Lee JY, Rein S. Strategies for improving the efficiency of Cz-silicon 

solar cells. In Conference Record, 28th IEEE Photovoltaic Specialists Conference. 

Anchorage September 2000; 201-204. 

[4-12] Lee JY, Peters S, Rein S, Glunz SW. Improvement of charge minority-carrier 

lifetime in p(boron)-type Czochralski silicon by rapid thermal annealing. Progress 

in Photovoltaics 2001; 9: 417-424. 

[4-13] Bothe K, Schmidt J, Hezel R. Effective reduction of the metastable Defect 

concentration in boron-doped Czochralski silicon for solar cells”, In Conference 

Record, 29th IEEE Photovoltaic Specialists Conference. New Orleans May 2002; 

194-197. 

[4-14] Crabtree G, Jester TL, Fredric C, Nickerson J, Meemongkolkiat V, Rohatgi A. 

Production viability of gallium doped mono-crystalline solar cells. In Conference 

Record, 31st IEEE Photovoltaic Specialists Conference. Lake Buena Vista January 

2005; 935-938. 

[4-15] Metz A, Abe T, Hezel R. Gallium-doped Czochralski grown silicon: a novel 

promising material for the PV-industry. In Conference Proceedings, 16th European 

Photovoltaic Solar Energy Conference. Glasgow, May 2000; 1189-1192. 

[4-16] Diez S, et al. Alternatives to boron-doped Czochralski for silicon solar cell 

processing. In Conference Proceedings, 19th European Photovoltaic Solar Energy 

Conference. Paris, June 2004; 504-507. 

[4-17] Münzer KA, Holdermann KT, Schlosser RE, Sterk S. Thin monocrystalline 

silicon solar cells. IEEE Transactions on Electron Devices 1999; ED-46: 2055-

2061. 

[4-18] Damiani B, Ristow A, Ebong A, Rohatgi A. Design optimization for higher 

stabilized efficiency and reduced light-induced degradation in boron-doped 

Czochralski silicon solar cells. Progress in Photovoltaics 2002; 10: 185-193. 

[4-19] Sinton RA, Cuevas A. Contactless determination of current-voltage characteristics 

and minority-carrier lifetimes in semiconductors from quasi-steady-state 

photoconductance data. Applied Physics Letters 1996; 69: 2510-2512. 

[4-20] Nagel H, Berge C, Aberle AG. Generalized analysis of quasi-steady-state and 

quasi-transient measurements of carrier lifetimes in semiconductors. Journal of 

Applied Physics 1999; 86: 6218-6221. 

[4-21] Fan TW, Qian JJ, Wu J, Lin LY, Yuan J. Tentative analysis of Swirl defects in 

silicon crystals. Journal of Crystal Growth 2000; 213: 276-282. 

[4-22] Polignano ML, Cerofolini GF, Bender H, Claeys C. Gettering mechanisms in 

silicon. Journal of Applied Physics 1988; 64: 869-876. 

[4-23] Glunz SW, Rein S, Lee JY, Warta W. Minority carrier lifetime degradation in 

boron-doped Czochralski silicon. Journal of Applied Physics 2001; 90: 2397-

2404. 

[4-24] Clugston DA, Basore PA. PC1D version 5: 32-bit solar cell modeling on personal 

computers. In Conference Record, 26th IEEE Photovoltaic Specialists Conference. 

Anaheim September/October 1997; 207-210. 

[4-25] Cuevas A, Stocks M, Macdonald D, Kerr M, Samundsett C. Recombination and 

trapping in multicrystalline silicon. IEEE Transactions on Electron Devices 1999; 

ED-46: 2026-2034. 

56
5. Understanding the Formation of High-Quality Thick-Film Ag 

Contacts on High Sheet-Resistance Si Emitters for Solar Cells 

5.1 Introduction 

Screen-printed solar cells have been around for about three decades; however, little is 

understood about the physics and chemistry of the formation of thick-film silver contacts 

to the Si emitter, as well as carrier transport at and near the Ag-Si interface and its precise 

correlation with the quality of the ohmic contact. Reasonably good quality screen-printed 

Ag contacts are routinely achieved on 45 Ω/sq emitters for solar cells. However, making 

good ohmic contacts to high sheet-resistance (80-100 Ω/sq) emitters is still a challenge. 

Several manufacturers around the world are working and partly succeeding on the 

development of Ag pastes that work for high sheet-resistance emitters [5-1]. Silver thickfilm pastes typically contain silver powder, glass frit, and organic binder. The glass frit is 

important in influencing the contact quality because it can affect adhesion, conduction, and 

contact formation. The glass frit normally consists of lead oxide containing silicate glass. It 

has been shown that Ag dissolution in the glass frit occurs upon firing [5-2]. Recently, Ag 

crystallites have been shown to grow into the Si surface from the glass [5-3, 5-4]. The lead 

borosilicate glass layer over most of the interface is assumed to be insulating because of its 

high resistivity. Thus, current transport is projected to take place via islands of Ag and Si 

contact [5-5]. Ballif et al. validated the presence of Ag crystallites at the Ag-Si interface 

along with very low contact resistivity (~2×10-7 Ω-cm

2

) between these Ag crystallites and 

the Si emitter coupled with the very high resistivity (109 Ω-cm) of glass and thus current 

transport from the Ag crystallites to the Ag bulk through ultrathin glass regions has been 

proposed [5-6]. It has also been proposed [5-4] that conduction takes place via multi-step 

tunneling via metal precipitates in the glass layer between Ag crystallites at the Ag-Si 

interface and the Ag bulk of the grid. Moreover, the chemistry of glass frit/modifiers can 

greatly influence junction shunting and the junction leakage current because of etching of 

Si by the glass frit during the contact firing cycle and possible migration of metal 

impurities present in the glass into the emitter region [5-7- 5-9]. The above ideas have been 

used and expanded to explain why, of the three Ag pastes we have studied in this paper, 

PV168 paste from Dupont gave high fill factors (FFs) while the other two pastes did not 

work as well on 100 Ω/sq emitters [5-10, 5-11]. To gain a better understanding of the 

requirements for achieving good screen-printed ohmic contacts, two other widely used 

commercial thick-film pastes (A and B) were studied in an attempt to analyze and correlate 

the electrical and physical behaviors of screen-printed contacts. 

5.2 Experimental Method 

In this study, screen-printed n+-p-p+ solar cells (4 cm2) were fabricated on single crystal 

Si using different Ag pastes and firing conditions on a 100 Ω/sq emitter. P-type, 0.6 Ω-cm, 

300-μm-thick (100) float-zone (FZ) substrates were used for all the experiments. Silicon 

wafers were first chemically cleaned, followed by POCl3 diffusion to form the n+

-emitter. 

After the phosphorus-glass removal and another clean, plasma-enhanced chemical vapor 

57
deposited (PECVD) SiNx antireflection (AR) coating was deposited on the emitter. Next, 

an Al paste was screen-printed on the backside and dried at 200°C. A Ag grid was then 

screen-printed on top of the SiNx film, and Ag and Al contacts were co-fired (single firing 

step) in a lamp-heated three-zone belt-line furnace. All the cells were then isolated using a 

dicing saw and were annealed in forming gas at 400°C for ~15 minutes. Cross-sectional 

scanning electron microscopy (SEM) and plane-view atomic force microscopy (AFM) 

measurements were taken to study the structure of the Ag-Si interface. Cross-sectional 

conductive AFM (CAFM) images were taken to investigate the conductivity through the 

bulk of the PV168 Ag gridline for high and low firing temperatures. The CAFM technique 

uses a bias voltage applied between a highly conductive tip and the Al back contact as the 

tip passes over the cross-sectional grid. No etching was found necessary for the crosssectional SEM and CAFM measurements. However, the Ag grid was etched prior to the 

AFM using HF, HNO3 and HF acids in sequence to remove the metal and the glass to 

delineate the footprints of the etched Ag crystallites by removing the metal and the glass. 

The plane-view AFM images of the surface morphology were acquired in intermittent 

tapping mode. Cross-sectional TEM was performed on selected samples. The sample 

preparation for the TEM measurement involved standard mechanical polishing down to 

about 50 μm followed by dimple polishing down to a thickness around 5μm. The samples 

were then ion beam milled using a Gatan precision ion polishing system (PIPS) at low 

energy, and low angle. Secondary-ion mass spectroscopy (SIMS) measurements were 

taken on selected samples to determine the Ag and P concentration profiles in the silicon 

directly underneath the Ag grid using a CAMECA IMS-5F ion microscope. The Ag grid 

was etched in HCl for the SIMS measurements to prevent any Si etching. The samples 

were bombarded with 5.5 KeV O2

+

 primary ions for the analysis of Ag whereas14.5 KeV 

Cs+

 primary ions were used for P. The primary ion beam was rastered over an area of 150 

μm × 150 μm. Positive or negative secondary ions were collected from a ~20-μm-diameter 

area in the center of the rastered crater for Ag and P, respectively. Contact resistance was 

determined by using the transfer-length method (TLM) pattern [5-12], printed 

simultaneously with the grid, while the cell parameters were measured by dark and light 

I-V measurements. 

5.3 Results and Discussion 

Three different screen-printing pastes (A, B, and PV168) were investigated using SEM, 

AFM, and SIMS analysis in conjunction with contact resistance and solar cell performance 

measurements to understand and establish a correlation between physical and electrical 

properties of screen-printed contacts on a high sheet-resistance emitter (100 Ω/sq). Two of 

the pastes (A and B) are widely used in industry, and the third paste (PV168) is 

commercially available from Dupont, but its use is not as widespread. All the three pastes 

contained some P which could result in self-doping. Therefore, SIMS measurements were 

first taken to check the formation of selective emitter via the injection of P underneath the 

gridlines. This was determined by printing and firing the three pastes on an undiffused 

wafer (no emitter). Figure 5-1 shows that all three pastes introduced limited amount of P 

near the Si; thus, for the high-temperature (835°C) firing condition used in this study 

injected P profiles resulted in sheet resistance of 900-5000 Ω/sq. More important, the 

58
1.00E+16

1.00E+17

1.00E+18

1.00E+19

1.00E+20

1.00E+21

0.00 0.05 0.10 0.15 0.20 0.25 0.30

Depth (μm)

P Concentration from the Paste Only (cm-3)

Paste A

Paste B

PV168

P 100 Ohm/sq emitter profile

self-doping P concentration is much smaller than the P concentration associated with the 

diffusion of 100 Ω/sq emitter (Fig. 5-1). Since the injected P concentration does not alter 

the net P concentration underneath the grid, selective emitter formation can be precluded 

(unless it is within a few atomic layers, which may be affected or dissolved during the 

removal of Ag with HCl). Figure 5-1 also shows that the PV168 paste introduces much 

less P into Si compared to pastes A or B when fired at high temperatures. Thus, we can 

ignore self-doping as the basis for any similarity and differences between the contacts 

formed by these pastes. This led to the SEM and AFM investigation of Ag crystallites and 

interface structure in the following sections to explain the difference in the electrical 

performance of the contacts. 

Figure 5-1. SIMS profiles of P detected in Si after 835°C fast firing of pastes A, B, and 

PV168 on undiffused Si. P profile for a 100 Ω/sq emitter is also shown for comparison. 

5.3.1 Effect of Firing Temperature on the Ag-Si Contact Interface for PV168 Paste 

Table 5-1 summarizes the electrical performance of contacts and solar cells fabricated 

with PV168 Ag paste on 100 Ω/sq emitter using conventional 750°C firing and a highertemperature (835°C) firing. It is clear from Table 5-1 that the lower-temperature (750°C) 

co-firing of PV168 paste results in unacceptably high contact resistance (45 mΩ-cm

2

) in 

conjunction with low FF (0.643) and poor cell efficiency (14.37%). However, the 835°C 

firing produced excellent ohmic contacts with specific contact-resistance of ~2 mΩ-cm

2

and series resistance of 0.7 Ω-cm2

, resulting in a FF of 0.77 on a 100 Ω/sq emitter and a FZ 

Si cell efficiency of 17.4%. 

Table 5-1. The electrical performance of contacts and Si solar cells formed with paste 

PV168 on a 100 Ω/sq emitter using low- (conventional) and high-temperature firing. 

59
To gain a better understanding of the cell results, contact interface analysis was 

performed using SEM, AFM, and SIMS. The SEM image (Fig. 5-2(a)) of the interface of 

the Ag-Si contact fired at 835°C revealed the presence 50-200-nm wide re-grown Ag 

crystallites. Figure 5-2(b) shows the AFM image of the plane view of the same sample (for 

835°C firing) after etching the Ag and the glass frit. Distributed pits observed in the AFM 

image are indicative of the footprints of the Ag crystallites, which were removed during 

the Ag etching prior to AFM. The AFM image in Figure 5-2(b) also reveals that the depth 

of these crystallites in the Si is on the order of 50 nm (notice that the depth scale goes from 

0-100 nm), which is much smaller than the junction depth of ~0.28 μm for the 100 Ω/sq 

emitter. This is in good agreement with the slight shift observed in the SIMS P profiles for 

regions with and without the gridline, which showed a displacement of ~55 nm, 

approximately equal to the penetration depth of the Ag crystallites for the PV168 paste [5-

11]. The Ag crystallite penetration depth of only ~50 nm also explains why there was no 

appreciable shunting in the case of PV168 paste even after the 835°C firing. The Rsh and 

Jo2 values were very reasonable (131.4 kΩ-cm

2

 and 15 nA/cm2, respectively), contributing 

to the observed high FF of 0.782 (Table 5-1). 

The AFM images in Figures 5-2(b) also reveals a larger number of Ag crystallites than 

reported in [6] for a conventional paste and firing scheme. The regular transfer of dissolved 

Ag through the glassy layer to the Si surface, resulting in a regular distribution of a large 

number of Ag crystallites (shown in Fig. 5-2(a) and 5-2(b)) could be attributed to the lower 

glass transition temperature which has been shown to precipitate a larger amount of Ag 

crystallites [5-1]. The higher-temperature firing (835°C) dissolves more Ag into the glass 

frit, possibly reaching saturation, and upon cooling; the excess Ag may precipitate at the 

Ag-Si interface as crystallites. PV168 paste composition is such that even at ~835°C firing, 

which was necessary to give acceptable contact and series resistance, Ag crystallites were 

small enough (much less than the 0.28 μm junction depth) to avoid junction shunting. This 

was the key to achieving high FF on a 100 Ω/sq emitter with PV168 paste. For the lower 

temperature firing of 750°C Ag crystallites were very rare or generally not observed at all 

using high-magnification SEM. Hence, to explain this behavior TEM analysis was 

performed on the 750° C-fired sample. The TEM micrograph in Fig. 5-3 shows that a 

SiNO layer, which has a variable thickness, is not completely etched by the glass frit 

indicating a less aggressive glass frit chemistry compared with conventional pastes, which 

achieve good ohmic contacts at conventional firing temperatures. The thickness of the lead 

borosilicate glass varies extensively as shown in Fig. 5-3. The thickness of the SiNO layer 

is ~65 nm which is too thick for tunneling to take place. Full penetration of the 

antireflection coating is necessary for a good contact with the emitter leading to low 

contact resistance [5-7]. This explains the high contact resistance observed for the PV168 

Ag paste when fired at conventional firing temperatures of ~750°C. Regions where the 

SiNO is completely etched could be very sparse at this firing temperature. 

Over-firing increases the aggressiveness of the etching reaction between the glass frit 

and the SiNO layer. Thus, the SiNO is completely etched (in most regions), allowing the 

glass frit to reach and dissolve the underlying Si which recrystallizes upon cooling with the 

Ag crystallites growing epitaxially as shown in Fig. 5-2(a). 

60
Figure 5-2(a). SEM image of the Ag-Si interface for PV168 fired at 835°C. 

Figure 5-2(b). AFM plane-view of the Si interface for PV168 fired at 835°C. 

Figure 5-3. TEM micrograph of a 750°C-fired PV168 Ag/Si contact. 

61
Figures 5-4(a) and 5-4(b) show the cross-sectional conductive AFM images of the bulk 

of the Ag gridline and the Si beneath it for 750°C and 835°C firing temperatures. This 

technique can probe the variation in the conductivity of the gridline from top to bottom. 

The presence of a glass layer in the grid or at the interface can lower the grid conductivity 

in those regions. The bright contrast shows high conductivity or high current between the 

probes for the same probing voltage applied between the back Al and the probing point in 

the grid. CAFM conductivity results can also be influenced by (a) the conductivity of the 

grid Ag, which could be a function of Ag grain size and sintering, (b) the area fraction of 

the Ag crystallites at the interface, (c) the number of direct contact points between the bulk 

of the Ag grid and the Ag crystallites at the interface, and (d) the conductivity of the frit or 

tunneling through the glass layer (where it is thin enough) to complete the current path 

between Ag crystallites at the interface and the bulk of the grid. Although a few direct 

interconnections between the interface Ag crystallites and the Ag bulk might exist, none 

were observed using both the high magnification SEM and TEM. The observed nonhomogeneous conductivity of the Ag gridline in Figure 5-4, particularly near the Ag-Si 

interface, is attributed to the presence of varying thickness of high-resistivity glass between 

Si and the Ag grid. 

The CAFM measurements also confirmed that both firing temperatures gave nonrectifying ohmic contacts because similar current values were obtained when the applied 

voltage polarity was switched. However, for the lower firing temperature of 750°C, 

absolute conduction through the bulk of the PV168 gridline was lower (~350 nA current at 

200 mV) compared to 1800 nA current at 200 mV. To achieve the same current (≥1800 

nA) for the 750°C-fired grid, the voltage applied to the grid had to be increased 

significantly from 200 mV to 1000 mV. 

The Ag particle size, shape, surface area, surfacant treatment, as well as the glass frit all 

affect the sintering behavior of the Ag gridline [5-13,5-14]. The characteristics of the Ag 

particles and glass frit in PV168 are such that the Ag particles do not sinter as well at 

750°C to produce higher Ag grain size within the bulk of the grid. It has been shown in 

[5-15] that the Ag grain growth increases with firing temperature, yielding denser films 

and causing the sheet resistivity of the Ag thick films to decrease with higher sintering 

temperatures in the range of 450°C -800°C. We found similar behavior for the PV168 Ag 

paste (Figs. 5-5(a) and 5-5(b)) where Ag grain size was of ~4.5 μm after the 835°C firing 

as opposed to ~2.1 μm for the 750°C firing. In addition, the grains formed at 835°C firing 

showed a very compact structure. This yielded a lower gridline resistivity of ~1.9 μΩ-cm, 

which is close to the resistivity of pure Ag (~1.7 μΩ-cm). The gridline resistivity after the 

750°C firing was ~2.5 μΩ-cm. 

62
Voltage = +200 mV 

10 μm

Figure 5-4(a). Cross-sectional conductive AFM of the Ag gridline fired at 750°C. 

Si Ag 

Ag-Si 

Interface 

Voltage = +200 mV 

Si Ag 

15 μm 

Ag-Si 

Interface 

Figure 5-4(b). Cross-sectional conductive AFM of the Ag gridline fired at 835°C. 

63
Figure 5-3 shows that 835°C firing produces larger Ag crystallites with increased Ag-Si 

contact fraction. If current transport mostly takes place via tunneling through ultra thin 

glass regions as proposed in [5-6], an increase in the area coverage of the Si emitter by the 

Ag crystallites would also increase the probability of encounter of these regions. Also, if 

the current transport takes place via direct interconnection between the Ag bulk and 

interface crystallites, an increase in the Ag-Si area coverage due to a larger number of Ag 

crystallites would also decrease the contact resistance. Finally, the relative conductivity of 

the darker regions near the Ag-Si interface in Figure 5-4, may be indicative of glass 

conductivity which is lower than that of the Ag. The electrical conductivity model for 

discontinuous metal structures is based on the tunneling process between metal granules 

[5-16]. Pb atoms usually agglomerate into large Pb crystallites, and thus, tunneling is 

difficult to occur; Bi ions have been found to decrease the conductivity of lead-silicate 

glasses significantly because they act as tunneling centers [5-16]. Bi has been observed in 

the glass frit in PV168 Ag paste (Fig. 5-3), which would explain why the glass is 

conductive, although not as well as Ag, as shown in Fig. 5-4. The conductivity is lower for 

both the Ag and glass regions for the 750°C firing due to the incomplete etching of the 

SiNO layer. Thus, the effective series resistance for the 835°C firing is only 0.854 Ω-cm

2

as opposed to 4.39 Ω-cm

2

 for the 750°C firing (Table 5-1). The improvement in the 

sintering of the Ag particles for the higher firing temperature (835°C) is not significant 

enough to yield the observed improvement in conductivity shown in the CAFM images in 

Figure 5-4(b). From Figs 5-3 and 5-4 limited current transport between the Si emitter and 

the Ag bulk can be attributed as the main reason for the high contact resistance for the 

750°C fired cell because of the incomplete etching of the dielectric layer. By increasing the 

aggressiveness of the glass frit in PV168 the SiNO layer could be completely etched at 

750°C. Regions where the glass is very thin (Fig. 5-3) could help improve the current 

transport via tunneling or by facilitating the existence of a direct connection between the 

Ag crystallites and Ag bulk. 

To understand the role of the frit we investigated another DuPont paste (PV167), which 

is analogous to PV168 but without the glass frit. We did not get proper adhesion with 

PV167 paste when fired at 835°C for a few seconds. This suggests that for the fast firing 

condition, the presence of the glass frit plays an important role in achieving proper contact 

even at the Ag-Si eutectic temperature of 835°C [5-17]. Thus, the glass frit helps to 

adequately dissolve Ag at the Ag-Si eutectic temperature to produce large uniformly 

distributed crystallites. In addition, the Ag also sinters well at 835°C to give large Ag 

grains. Even though the PV168 Ag paste was initially designed to operate as a self-doping 

paste at temperatures greater than the Ag-Si eutectic of 835°C, it appears that the Ag-Si 

alloying temperature is also significant in that it is the temperature where the glass frit and 

Ag particles of PV168 were designed to operate to give better contact quality [5-18]. This 

is why the PV168 performs well around the Ag-Si eutectic temperature and fails at the low 

firing temperature for high sheet-resistance emitters. 

64
5.3.2 Investigation of Screen-Printed Contacts to 100 Ω/sq Emitter Using Widely Used 

Commercial Pastes A and B and Conventional Firing Temperature 

After establishing that PV168 can form good contacts to a 100 Ω/sq emitter at 835°C, 

we investigated the performance of two widely used Ag pastes A and B. Commercial 

pastes are generally fired at sample temperatures around 750°C to avoid junction shunting. 

This process works well for the 45 Ω/sq emitter. Table 5-2 shows that pastes A and B gave 

reasonably low Rs (~0.8 Ω-cm

2

), Jo2 (~25 nA/cm2), and high Rsh (>7 kΩ-cm

2

) on 45 Ω/sq 

emitter, producing FFs of greater than 0.77 and cell efficiency >16.7% on FZ Si. However, 

Table 5-3 shows that both pastes failed on the 100 Ω/sq emitter with the 750°C firing 

condition. 

Table 5-2. Electrical performance of contacts and solar cells formed with Ag pastes A and B 

using 45 Ω/sq emitters and conventional firing temperatures of 750°C 

Table 5-3 Electrical performance of contacts and solar cells formed with Ag pastes A and B 

he plane-view AFM photograph in Figure 5-5(a) shows that the 750°C firing of paste 

A p

gure 5-5(b) shows that paste B behaves quite differently than 

pas

using 100 Ω/sq emitters and conventional firing temperature of 750°C. 

T

roduced a highly irregular distribution of Ag crystallites. In addition, the contact area is 

quite small, which may not be suitable for high sheet-resistance emitters. This observation 

is consistent with reference [5-6] for conventional paste and firing conditions and explains 

the higher specific contact-resistance of ~4 mΩ-cm2

 on 100 Ω/sq emitter (Table 5-3). 

Table 5-2 shows that the same firing condition produced a contact resistance of 0.24 mΩcm

2 on the 45 Ω/sq emitter, indicating the need for more or larger Ag crystallites for high 

sheet-resistance emitters. 

The AFM image in Fi

te A for the 750°C firing. Paste B showed a fairly uniform distribution of small Ag 

crystallites (~25 nm in width). A regular distribution of small Ag crystallites is better for 

lower contact resistance compared to an irregular distribution of larger Ag crystallites, 

which are limited to certain regions only. The SEM image in Fig. 5-7 reveals that the Ag 

crystallites formed at the Ag-Si interface for paste B are on the order of 25-120 nm in size; 

thus, in some cases they are as large as those formed in the case of PV168 when fired at a 

higher temperature of 835°C. However, paste B consists of a more aggressive glass frit 

than PV168 Ag paste, which leads to excessive Ag penetration into the junction. This was 

confirmed by the SIMS measurements in Fig. 5-8, which reveal the presence of ~3×1017

65
cm-3 Ag at or near the n+

-p junction which is ~0.28 μm deep. This increased the leakage 

current to 2678 nA/cm2 (Table 3) and gave a lower FF of 0.635. 

 has been demonstrated [5-19 – 5-21] that, in addition to promoting adhesion, proper 

gla

Thus, the led on the 100 

Ω/sq emitter at the conventional firing temperature of 750°C. Paste A failed because of 

hig

It

ss frit chemistry can actually retard the Ag migration into the emitter, reducing 

shunting. Unlike the glass frit in paste PV168, the glass frit in paste B is not effective in 

arresting Ag close to the Si surface to provide good-quality contact to the shallow high 

sheet-resistance emitter even at 750°C firing. 

above analysis explains why pastes A, B and PV168 all fai

her contact resistance, attributed to non-uniform distribution of Ag crystallites and low 

contact area fraction; paste B failed because of the high junction leakage associated with 

excessive Ag and impurity penetration into the emitter; PV168 failed because of the sparse 

distribution of smaller Ag crystallites and low contact area fraction, resulting in high 

contact resistance. 

Figure 5-5. AFM plane-view of the Si interfac a) paste A and (b) paste B both fired at e for (

(a) (b)

750°C. 

Figure 5-7. SEM of the Ag-Si interface for Paste B fired at 750°C. 

66
1.00E+16

1.00E+17

1.00E+18

1.00E+19

1.00E+20

1.00E+21

0.00 0.10 0.20 0.30 0.40

Depth (μm)

Concentration (cm -3)

Ag P

Juncti

on

Figure 5-8. SIMS Ag and P profiles in Si for paste 

B at 750°C firing. 

5.3.3 High Temperature (835°C) Firing of Conventional Pastes A and B on the 100 Ω/sq 

Emitter 

The final step in applying this methodology to explain the performance of screenprinted contacts involved the study of the contact interface resulting from 835°C firing of 

pastes A and B. This is because PV168 worked at high temperature even though it failed at 

low temperature. The solar cell and contact data in Table 5-4 show that both the pastes 

failed again at 835°C, with the ideality factor (n) rising above three in conjunction with 

very high leakage current exceeding 14,000 nA/cm2

. At 835°C, paste A gave very good 

contact resistance of 0.59 mΩ-cm

2

 in conjunction with series resistance of 0.83 Ω-cm

2

, but 

it failed because of a very high junction leakage current (Jo2). The EDS measurements (Fig. 

5-9) detected the presence of several metal elements in the glass frit of paste A at a very 

high concentration, deep into the emitter region. In addition, the SIMS measurements (Fig. 

5-11) showed the presence of high Ag concentration at the junction. 

Paste B failed again because of excessive shunting after the 835°C firing. The SEM 

image in Figure 5-10 shows a significant increase in the size of Ag crystallites, 

approaching 0.6-1 μm. Since this is larger than the junction depth of ~0.28 μm, it explains 

why the 835°C-fired paste B cells were so badly shunted. A comparison of Figs. 5-7 and 5-

10 show that the Ag crystallites grew almost 10 times in size (~100 nm⇒ ~1000 nm) when 

the paste B firing temperature was raised from 750°C to 835°C. This is probably because 

the glass frit in paste B was able to dissolve a lot more Ag at the higher temperature and 

became highly supersaturated, producing large precipitates or crystallites at the interface 

upon cooling. This is also consistent with the Ag-Pb phase diagram [5-22]. The Ag SIMS 

profiles in Figure 5-11 clearly show that pastes A and B failed because of Ag-induced 

junction shunting while PV168 survived 835°C firing because of the significantly reduced 

penetration of Ag into the 100 Ω/sq emitter. 

67
Table 5-4. Electrical performance of solar cells made with pastes A and B on 100 Ω/sq 

emitters using high-temperature firing. 

1.00E+20

1.01E+22

2.01E+22

3.01E+22

4.01E+22

5.01E+22

6.01E+22

Ag glass

element 1

glass

element 2

glass

element 3

Si

Concentration (cm-3)

Figure 5-9. EDS analysis of the elements in the emitter 

region underneath the Ag contact for paste A. 

68
Figure 5-10. SEM images of the Ag-Si interface for paste B fired at 835°C. 

1.00E+15

1.00E+16

1.00E+17

1.00E+18

1.00E+19

1.00E+20

1.00E+21

0.0 0.1 0.2 0.3 0.4

Depth (μm)

Concentration (cm-3)

Paste B Ag

PV168 Ag

Paste A Ag

100 Ohm/sq P

Figure 5-11. SIMS analysis of the three pastes fired at 835°C. 

5.4 Conclusion 

A combination of SEM or TEM with EDS, and AFM are very useful for understanding 

the contact structure and composition. These techniques along with the high detection limit 

of the SIMS analyses has been used to understand and explain the electrical behavior of the 

screen-printed contacts to high- and low-sheet-resistance emitters. Glass frit chemistry, Ag 

particle size and shape, and surface treatment of Ag particles should be tailored to 

minimize Ag migration into the emitter while producing uniform distribution of Ag 

crystallites with sufficient contact area fraction, which can give good contact resistance 

without junction shunting. Higher firing temperatures produce larger Ag crystallites. 

Proper glass frit chemistry can allow for complete etching of the SiNx layer while arresting 

the diffusion of Ag into the junction as well. 

We found that all the three pastes (A, B, and PV168) investigated in this study failed on 

the 100 Ω/sq emitter at a firing temperature of ~750°C. Paste A failed because of the nonuniform distribution of Ag crystallites with small contact area fraction. Paste B failed 

69
because of excessive Ag penetration. PV168 failed because of low contact area fraction 

resulting from incomplete etching of the SiNx coating. At a firing temperature of 835°C, 

pastes A and B failed on 100 Ω/sq emitter because of junction shunting. Paste A failed 

because of an aggressive frit allowing excessive penetration of Ag along with other 

impurities into the emitter while paste B produced Ag crystallites larger than the junction 

depth. However, PV168 succeeded on the 100 Ω/sq emitter because the 835°C firing 

allowed the formation of uniform Ag crystallites of reasonable size, ~50 nm or much 

smaller than the junction depth of ~0.28 μm, while the frit chemistry maintained the Ag 

penetration below 0.2 μm. Thus, the combination of SEM, AFM, and SIMS measurements 

is quite effective in understanding the electrical behavior of screen-printed contacts and 

can lead to the development of screen-printed pastes that can produce high FF and cell 

performance at appropriate firing temperatures. 

5.5 References 

[5-1] M. M. Hilali, A. Rohatgi, C. Khadilkar, S. Kim, T. Pham, J. Salami, A. Shaikh, and 

S. Sridharan, “Development of Thick-Film Ag Paste to Achieve High Fill Factors 

in Screen-Printed Si Solar Cells with Low- and High-Sheet-Resistance Emitter,”

Proc. of 19th European Photovoltaic Solar Energy Conference, Paris, France, June 

2004, in press. 

[5-2] M. Prudenziati, L. Moro, B. Morten, and F. Sirotti, “Ag-Based Thick-Film Front 

Metallization of Silicon Solar Cells,” Active and Passive Electronic Components, 

Vol. 13, pp. 133-150. 

[8-3] C. Ballif, D. M. Huljić, A. Hessler-Wysser, and G. Willeke, “Nature of the Ag-Si 

Interface in Screen-Printed Contacts: A Detailed Transmission Electron 

Microscopy Study of Cross-Sectional Structures,” Proc. of the 29th IEEE PVSC, 

New Orleans, 2002, pp. 360-363. 

[8-4] G. Schubert, B. Fischer, and P. Fath, “Formation and Nature of Ag Thick Film 

Front Contacts on Crystalline Silicon Solar Cells,” PV in Europe Conference, 

Rome (2002). 

[8-5] K. Firor, S. J. Hogan, J. M. Barrett, and R. T. Coyle, “Series Resistance Associated 

with Thick-Film Contacts to Solar Cells,” Proc. of the 16th IEEE PVSC, 1982, pp. 

824-827. 

[8-6] C. Ballif, D. M. Huljić, G. Willeke, and A. Hessler-Wysser, “Silver Thick-Film 

Contacts on Highly Doped N-type Silicon Emitters: Structural and Electronic 

Properties of the Interface,” Appl. Phys. Lett., Vol. 82, No. 12, March 2003, pp. 

1878-1880. 

[8-7] R. J. S. Young and Alan F. Carroll, “Advances in Front-side Thick Film 

Metallizations for Silicon Solar Cells,” Proceedings of the 16th European 

Photovoltaic Energy Conference, Glasgow, UK, May 2000, pp. 1731-1734. 

[8-8] G. C. Cheek, R. P. Mertens, R. Van Overstraeten, and L. Frisson, “Thick-film 

Metallization for Solar Cell Applications,” IEEE Transactions on Electron Devices, 

Vol. ED-31, May 1984, pp. 602-609. 

70
[8-9] R. Mertens, M. Eyckmans, G. Cheek, M. Honore, and R. Van Overstraten, “Critical 

Processing Parameter Optimization for Screen-Printed Semicrystalline Silicon 

Solar Cells,” Proc. of the 17th IEEE PVSC, 1984, pp.1347-1351. 

[8-10] M. Hilali, J.–W. Jeong, A. Rohatgi, D. L. Meier, and A. F. Carroll, “Optimization 

of Self-Doping Ag Paste Firing to Achieve High Fill Factors on Screen-Printed 

Silicon Solar Cells with a 100 Ω/sq. Emitter,” Proc. of the 29th IEEE PVSC, New 

Orleans, May 2002, pp.356-359. 

[8-11] M. M. Hilali, A. Rohatgi, and S. Asher, “Development of Screen-Printed Silicon 

Solar Cells with High Fill Factors on 100 Ohms/sq Emitters,” IEEE Trans. On 

Elect. Dev., Vol. 51, No. 6, June 2004, pp. 948-955. 

[8-12] D. K. Schroder, Semiconductor Material and Device Characterization, New York: 

John Wiley & Sons, Inc., 1990. 

[8-13] J. C. Lin, and C. Y. Wang, “Effect of Surfacant Treatment of Silver Powder on the 

Rheology of its Thick-Film Paste,” Materials Chemistry and Physics, Vol. 45, 

1996, pp. 136-144. 

[8-14] S. B. Rane, P. K. Khanna, T. Seth, G. J. Phatak, D. P. Amalnerkar, and B. K. Das, 

“Firing and Processing Effects on Microstructure of Fritted Silver Thick Film 

Electrode Materials for Solar Cells,” Materials Chemistry and Physics, Vol. 82, 

2003, pp. 237-245. 

[8-15] J. C. Lin, and C. Y. Wang, “Effect of Surface Properties of Silver Powder on the 

Sintering of its Thick-Film Conductor,” Materials Chemistry and Physics, Vol. 45, 

1996, pp. 253-261. 

[8-16] O Growski, L. Murawski, and K. Trzebiatowsky, “The Surface Conductivity of 

Lead Glasses,” Journal of Physics D: Applied Physics, Vol. 15, 1982, pp. 1097-

1101. edited by T. B. Massalski, (American Society for Metals, Materials Park, 

OH, 1990) pp. 92-94. 

[8-18] D. L. Meier, Hubert P. Davis, Ruth A. Garcia, and Joyce A. Jessup, “Method and 

Apparatus for Self-Doping Contacts to a Semiconductor,” Ebara Corporation, U.S. 

patent #6,703,295, March, 2004. 

[8-19] C. Khadilkar, S. Kim, T. Pham, A. Shaikh, and S. Sridharan, “Characterization of 

Silver Front Contact in a Silicon Solar Cell,” Technical Digest of the14th 

International Photovoltaic Specialists Energy Conference, Bangkok, Thailand, 

2004, pp. 443-444. 

[8-20] A. Shaikh, S. Sridharan, T. Pham, and C. Khadilkar, “Designing a Front Contact 

Ink for SiNx Coated Si Solar Cells” ” 3rd World Conference on Photovoltaic Energy 

Conversion, Osaka, Japan, May, 2003. 

[8-21] S. Sridharan, C. Khadilkar, T. Pham, and A. Shaikh “Characterization of 

Silver/Glasss/Silicon Front Contact Interface in a Silicon Solar Cell” presented at 

the 13th Workshop on Crystalline Silicon Solar Cell Materials and Processes, Vail, 

Colorado, August, 2003, pp. 162-165. 

[8-22] I. Karakaya and W. T. Thompson in Binary Alloy Phase Diagrams, 2nd edition 

edited by T. B. Massalski, (American Society for Metals, Materials Park, 

OH,1990), pp.72-73. 

71
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PLEASE DO NOT RETURN YOUR FORM TO THE ABOVE ORGANIZATION.

1. REPORT DATE (DD-MM-YYYY)

November 2007 

2. REPORT TYPE

Subcontract Report 

3. DATES COVERED (From - To)

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5a. CONTRACT NUMBER

DE-AC36-99-GO10337 

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4. TITLE AND SUBTITLE

Fundamental Research and Development for Improved Crystalline 

Silicon Solar Cells: Final Subcontract Report, March 2002 – 

July 2006 

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6. AUTHOR(S)

 A. Rohatgi 

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Georgia Institute of Technology 

Atlanta, Georgia 30332-0420 

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14. ABSTRACT (Maximum 200 Words)

This report summarizes the progress made by Georgia Tech in the 2002-2006 period toward high-efficiency, low-cost 

crystalline silicon solar cells. This program emphasize fundamental and applied research on commercial substrates 

and manufacturable technologies. A combination of material characterization, device modeling, technology 

development, and complete cell fabrication were used to accomplish the goals of this program. This report is divided 

into five sections that summarize our work on i) PECVD SiN-induced defect passivation (Sections 1 and 2); ii) the 

effect of material inhomogeneity on the performance of mc-Si solar cells (Section 3); iii) a comparison of light-induced 

degradation in commercially grown Ga- and B-doped Czochralski Si ingots (Section 4); and iv) the understanding of 

the formation of high-quality thick-film Ag contacts on high sheet-resistance emitters (Section 5). 

15. SUBJECT TERMS

PV; silicon; research and development; solar cells; crystalline; high efficiency; low cost; characterization; device 

modeling; defect passivation; light-induced degradation; Ag contacts; 

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