Microcrystalline materials and cells deposited by RF glow discharge

Solar Energy Materials & Solar Cells 78 (2003) 543–566

Microcrystalline materials and cells deposited by RF glow discharge Michio Kondo National Institute for Advanced Industrial Science and Technology, Umezono, Tsukuba, Ibaraki 305-0035 Japan

Abstract Recent developments in depositing high quality intrinsic and doped microcrystalline Si at low temperatures, 100–1401C, and high rates of B6 nm/s are reviewed. A new high-pressure depletion deposition method is described that suppresses ion bombardment and yields low defect densities. Passivation of oxygen related donors by hydrogen is discussed as well as the dissociation of passivated B–H dopant atoms by annealing at 2001C. Interface damage effects on superstrate and substrate cells are identified. The influence of the microcrystalline texture formed spontaneously during growth on optimizing optical confinement has been studied. r 2002 Published by Elsevier Science B.V. Keywords: Microcrystalline; RF glow discharge

1. Introduction Microcrystalline silicon has been developed as a bottom cell in combination with a top cell made of hydrogenated amorphous silicon, and very recently an efficiency of 14.5% (Voc ¼ 1:41 V, Jsc ¼ 14:4 mA/cm2, and FF¼0.719) has been reported [1]. The high Jsc can be obtained by using a transparent interlayer between the top and bottom cells [2]. In tandem structures, the bottom cell requires a thickness of several microns to obtain the high current density because of the indirect band gap of crystalline silicon. This implies the technological significance for reducing the defect density and increasing the growth rate at a deposition temperature compatible with amorphous silicon processing. Plasma enhanced chemical vapor deposition (PECVD) has been widely employed not only for amorphous silicon but also for microcrystalline silicon since the E-mail address: [email protected] (M. Kondo). 0927-0248/03/$ - see front matter r 2002 Published by Elsevier Science B.V. doi:10.1016/S0927-0248(02)00451-8

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discovery of hydrogenated amorphous silicon [3]. Microcrystalline silicon was also obtained by plasma CVD from gaseous sources [4–7]. The advantage of plasma processing is the low-temperature synthesis of materials in a high temperature phase due to its high-energy electron and low gas temperature (Broom temperature), which avoids the thermal damage of the substrate. This is quite important for cost reduction of solar cells. It is believed that the cell performance is a monotonic function of grain size [8]. This empirical relation is accounted for in terms of the carrier recombination at the defective grain boundaries. The recent development of microcrystalline silicon reveals the exceptionally high efficiency, which is caused by hydrogen passivation of the grain boundaries thanks to the low-temperature processing. The material properties under the low temperature and high deposition rate and the development of its improvement will be reviewed.

2. Preparation of microcrystalline silicon PECVD has been successfully employed for the fabrication of amorphous and microcrystalline silicon since Spear and LeComber succeeded in n- and p-type doping of a-Si:H using monosilane as a source gas [3]. A typical RF-PECVD system is shown in Fig. 1. The plasma is generated between two parallel electrodes and the substrate is placed at the grounded electrode (anode). The RF power is fed to the other electrode (cathode) through the matching network and the blocking capacitor. For mc-Si:H deposition, monosilane, SiH4 diluted by hydrogen has been most commonly used and there exists a certain threshold value for dilution ratio, R ¼H2/SiH4 [9]. It has been pointed out that the atomic hydrogen on the growing surface gives rise to crystal growth at temperatures much lower than the melting

RF power supply

Matching Network

~

Source gas

Cathode Plasma

Gas flow

Substrate Anode

Heater

Vacuum Pump

Fig. 1. Schematic diagram of the PECVD reactor for microcrystalline silicon deposition.

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temperature. The atomic hydrogen is generated from a dissociation of silane, whereas it recombines with silane as H + SiH4 - H2 + SiH3. In order to increase the atomic hydrogen flux density on the growing surface, the partial pressure of silane should therefore be lowered, that is, high hydrogen dilution or silane depletion is necessary for microcrystalline growth. The silane depletion occurs in the high power regime where the consumption of silane exceeds the silane supply. Thus, the atomic hydrogen density is the most crucial factor for the deposition of mc-Si:H (Fig. 2). Another important factor is the control of ions. The ion energy is measured by a Langmuir probe and controlled by the bias voltage at the mesh and substrate [10]. The plasma potential, which is usually positive with respect to the substrate, and the sheath voltage between the plasma and the substrate, accelerate the positive ions + such as H+ 2 and SiH3 . The crystalline volume fraction is evaluated by the ratio of the intensities of Raman scattering signal at wave numbers of 520.5 cm1 and 480 cm1 for crystalline and amorphous phase, respectively. This Raman intensity ratio Ic =Ia will be used as a measure of the crystallinity. The Raman ratio Ic =Ia as a function of the sheath voltage Vsheath shows a clear correlation, that is, the higher the ion energy, the worse the crystallinity is. This result implies the requirement of a lower ion energy for the better crystallinity [11]. Device grade mc-Si:H is typically deposited at around 2001C as for a-Si:H and this availability of a low temperature process enables one to employ a variety of materials as a substrate such as glass and polymer. However, the material properties and the device structure influence the optimum processing temperatures as mentioned later. The standard excitation frequency of the plasma is 13.56 MHz, but recently higher frequencies have become popular because of better film properties and higher deposition rates [12–14]. In the VHF plasma, electron temperature and plasma

6

Ic/Ia (arb. units)

5 4 3 2 1 0

1

10 Vp-Vs (V)

100

Fig. 2. Influence of the sheath voltage upon the crystallinity of microcrystalline silicon. The crystallinity is obtained by the intensity ratio of the Raman peak for crystalline and amorphous silicon.

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potential decreases while the electron density increases because of the more efficient energy gain mechanism of electrons by the wave-riding effect [10]. Thus, the high electron density and the low electron temperature, that is, lower ion damage are suitable for microcrystalline silicon deposition. The microscopic mechanism of the formation of microcrystalline silicon has been extensively argued in terms of the etching model [15,16], chemical heating model [17], surface diffusion model [18] and subsurface reaction model [19]. Although the role of hydrogen is still controversial among these models, the importance of atomic hydrogen is accepted.

3. Impurity effects on material properties of microcrystalline silicon In contrast to amorphous silicon, microcrystalline silicon is more sensitive to impurities. It has been reported that oxygen at the grain boundaries is electrically active [20,21] and oxygen-related donors are expected to precipitate at grain boundary regions. If that is the case, the accumulated donors at grain boundaries can cause a shunt leakage along the columnar structure, which is commonly observed in mc-Si:H. This shunt leakage results in lower open-circuit voltage Voc and fill factor FF. A possible way to avoid the impurity effect is to prepare purified material without oxygen using a ultrahigh vacuum system [22], although that is not applicable for actual solar cell manufacturing. An important findings, however, is the passivation of oxygen donors at low deposition temperature Ts o1801C. The Ts dependences of the carrier density n and mobility m in non-doped samples prepared on glass substrates are shown in Fig. 3. The carrier density is almost constant at Ts between 1001C and 1801C and increases by 4 orders of magnitudes at 2501C. The oxygen concentration of the samples was measured by secondary ion mass spectroscopy (SIMS) to be almost constant at

1017

1

Mobility (cm2/ V.s)

1015 1014 1013

0.1

1012 1011

Carrier density (cm-3)

1016

1010 0.01 50

100 150 200 250 i-layer deposition temperature (°C)

109 300

Fig. 3. Deposition temperature dependence of the mobility and the carrier density for undoped microcrystalline silicon. The career conduction was n-type for all the samples.

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2  1019 cm3 for all Ts : This result indicates that the reduction in n arises from the passivation of oxygen-related donors or from the compensation of donors by deep level defects such as dangling bonds. Fig. 4 shows the i-layer properties, such as the dependence of n; defect density N and hydrogen content CH on both Ts and annealing temperature. The defect density was measured by electron spin resonance (ESR). The dangling bond density as a function of Ts is indicated by open squares. The annealing temperature dependence for the sample deposited at 1401C is indicated by solid circles. The ESR-spin density

Carrier density (cm-3)

1016 1015 1014 1013 1012 1011

(a) 1010

Defect density (cm-3)

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1016

(b) 1015

Hydrogen content (%)

16

(c)

14 12 10 8 6 4 2 0

0

100

200

300

400

500

Temperature (°C) Fig. 4. Deposition temperature dependence of carrier concentration, defect density and hydrogen content of undoped samples for as deposited state (solid circles) and for as annealed state (open squares). The deposition temperature of annealed sample was 1401C.

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monotonically increases from 3.9  1015 to 3.9  1016 cm3 with decreasing Ts from 2501C to 1001C, as shown in Fig. 4(b). This Ts dependence of the ESR spin density is similar to the ones reported previously [12]. The decrease in N at 1801C as compared to 1401C is about 2–3  1015 cm3, while the increase in the n is several times greater and is 1  1016 cm3, suggesting an origin of increase in the n different from the compensation effect by the deep traps. Hence, the decrease in n can be explained in terms of the passivation of oxygen-related donors, probably by hydrogen. It has been proposed that three-fold coordinated oxygen forms a donor state in aSi:H [23]. Assuming a similar mechanism for oxygen donors incorporated at the grain boundary of mc-Si:H, hydrogen can passivate oxygen-related donors by H insertion into one of Si–O bond, i.e., O3Si +H-H–Si+Si–O–Si. Accordingly, Si–O–Si is no longer electrically active. The presence of such Si–H bonds adjacent to Si–O–Si bonds, however, has not been confirmed by IR-measurements. In order to confirm the stability of hydrogen passivation effects, we have performed thermal annealing at higher temperatures than Ts ¼1401C. With increasing annealing temperature up to 5001C, n remains constant as shown in Fig. 4(a). Nevertheless H decreases monotonically as shown in Fig. 4(c), accompanying an increase in the defect density. The most important result here is that the material deposited at temperatures lower than 1801C can maintain its intrinsic nature even after annealing at higher temperature. This will be utilized in solar cell processing. Crystal growth is often influenced by the presence of impurities. As shown in Fig. 5, B-doped microcrystalline silicon shows the best crystallinity at around 1801C. The best crystallinity is obtained at much higher temperatures for undoped films. For the B-doped layer, the hydrogen removal due to its catalytic reaction with boron on the growing surface deteriorates the crystallinity. This catalytic effect is consistent with an increase in the growth rate depending on the B doping ratio [24]. Lower Ts and higher hydrogen dilution, therefore, are required to obtain good crystallinity

15 Raman Crystallinity (lc/la)

B2H6/SiH4 = 0.03 % 14 13 12 11 10 80

100

120

140 160 180 Temperature [°C]

200

220

240

Fig. 5. Crystallinity measured by the Raman scattering for B-doped microcrystalline silicon as a function of the deposition temperature. Diborane concentration ratio to silane was 0.03%.

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particularly for very thin p-layers (o50 nm) without the loss of hydrogen coverage. However, a large amount of atomic hydrogen passivates the dopant atoms in the form of B–H complexes [25]. This complex dissociates into an electrically active center after 2001C annealing as reported in single crystalline silicon [26], and the solar cell performance is improved correspondingly. Fig. 6 shows the IR spectra of the mc-Si:H p-layer deposited at 1401C before and after 2001C thermal annealing. A peak near 1875 cm1 was found in the IR spectrum before the annealing as reported in the B-doped single crystalline silicon as a B–H complex [25]. The B–H passivates the B-acceptor because of the loss of the four-fold coordination. After annealing at 2001C, the peak near 1875 cm1 disappears and the carrier density increases due to the removal of H from the B–H complex [26]. In mcSi:H, the carrier density increases from 1018 to 1019 cm3 after thermal annealing at 2001C. Therefore, the carrier concentration in the p-layer can be increased by the thermal annealing. Thus, the combination of the low Ts and post annealing is more beneficial than high Ts deposition both for the intrinsic and doped layers. The presence of impurities also affects the grain growth of the crystallites. As shown in Fig. 7, the grain size is a monotonic function of the growth temperature. With decreasing Ts ; the grain size decreases gradually down to 1401C and decreases drastically below 1001C. The significant increase in defect density at 1001C shown in Fig. 4(b) is ascribed not only to the increase in defect density of the surrounding amorphous tissue but also to this rapid decrease in the grain size, because the volume fraction of defective grain boundary and the surrounding amorphous tissue increases. It has been reported that the purified material shows significant enhancement of the grain size in the temperature region higher than 3501C [22].

400 before annealing

350

after annealing 300

α

250 200

B-H

150 100 50 0

1800

2000 2200 Wave number (cm-1)

2400

Fig. 6. IR absorption spectra for B-doped microcrystalline silicon on c-Si substrates for as deposited sample (solid line) and annealed sample (dotted line) at 2001C.

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200 Grain size (Å)

150 1

100 50

100 150 200 250 i-layer deposition temperature (°C)

Orientation I (220) / I (111)

1.5

300

Fig. 7. The i-layer deposition temperature dependence of the crystalline grain size and the preferential orientation along [1 1 0] direction.

4. Low temperature approach of device grade microcrystalline silicon In the previous section, we demonstrated the advantage of the low temperature approach of microcrystalline silicon solar cells in spite of the disadvantages of defect density, crystalline volume fraction and grain size. The solar cell performance as a function of Ts for the pin type structure is shown in Fig. 8. The short-circuit current density Jsc seems to have a good correlation with the grain size as shown in Fig. 7, while the open-circuit voltage shows a complicated behavior. The fill factor has a good correlation with Voc (not shown). The reduction in Voc at higher Ts is ascribed to the leakage current along the grain boundary because of the oxygen related donors. The purified sample shows no significant decrease in Voc : The reduction in the lower temperature side may be explained in terms of an increase in the defect density that is an increase in the recombination current. Thus, the balance of Jsc and Voc ; the best efficiency Z ¼ 8:6% was obtained at 1401C [27]. This result was quite exceptional from a conventional point of view for polycrystalline silicon solar cells. In solar cell processing, the thermal annealing of the p-layer deposited at 1401C is important. When the p-layer is deposited at Ts >2001C, the Jsc decreases compared to the case of a p-layer deposited at 1401C with thermal annealing. This difference is ascribed to the band gap narrowing of the p-layer due to the catalytic role of boron to eliminate the surface hydrogen during growth [24]. Further decrease in the processing temperature has been studied by varying the hydrogen dilution condition, and as shown in Fig. 9, the efficiency can be maintained even at around 1001C, which enables us to fabricate microcrystalline silicon solar cells on plastic substrates as shown in Fig. 10(a). The I–V characteristics are shown in Fig. 10(b). The substrate used for this cell is a novel thermoplastic, ethylenetetracyclo-dodecene co-polymer (E/TD), which has a relatively high softening temperature of 1501C and a low humidity adsorption due to the absence of OH groups in the film (APELs Mitsui Chemical CO. Ltd.) [28].

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0.54

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Open Circuit Voltage (V)

Jsc (mA/cm2)

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0.44 250

200

Deposition Temperature (°C) Fig. 8. The i-layer deposition temperature dependence of short-circuit current and open-circuit voltage. Only the deposition temperature for the i-layer was varied and otherwise identical. The film thickness was fixed at 1 mm and the substrate used was Asahi-U type.

8

Efficiency (%)

6

4

2

0 50

100 Temperature (°C)

150

Fig. 9. The i-layer deposition temperature dependence of solar cell efficiency. The p- and n-layers were deposited at fixed conditions and the hydrogen dilution is optimized for each temperature.

5. Novel deposition technique for high throughput process Microcrystalline silicon solar cells require a thickness of several microns because of the indirect band gap. In particular, a recent approach for higher efficiency using a tandem structure requires even larger thicknesses than the single junction cell (3–5 mm). For mass production compatible with the a-Si:H top cell processing,

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(a)

25

Current Density (mA/cm2)

20

15

10

5

0 0 (b)

0.1

0.2

0.3

0.4

0.5

Voltage (V)

Fig. 10. (a) Picture of the microcrystalline silicon on a plastic substrate, and (b) its I–V characteristics.

a deposition time typically shorter than 3 min is needed, which implies a deposition rate over 5–10 nm/s. In plasma CVD, a high deposition rate requires high RF power, resulting in ion damage and a degradation of film quality [29,30]. As a novel approach for overcoming this dilemma, a high-pressure depletion (HPD) method has been proposed using a high deposition pressure above several

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6 8 Pressure (Torr)

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Raman Crystallinity (Ic/Ia)

Depo.Rate (A/s)

hundred Pa in combination with the depletion condition of the high RF power region [11,31,32]. The basic idea is that ion bombardment is suppressed under the high-pressure condition and that the atomic hydrogen density can be increased under depletion conditions. Although the generation rate of film precursors such as SiH3 is determined by the electron density as well as the molecule density, both the high pressure and high RF power are required for good quality materials. As shown in Fig. 11, the deposition rate increases with deposition pressure up to 4 Torr, while the crystalline volume fraction decreases monotonically. The subsequent decrease in the deposition rate is ascribed to the decrease in the electron temperature that is confirmed by the optical emission intensity from the plasma. The deterioration of the crystallinity is explained in terms of the atomic hydrogen density. The atomic hydrogen far from the substrate cannot reach the growing surface because of the recombination reaction. H + SiH4 - H2 + SiH3. Under the depletion condition, the SiH4 density and the recombination rate decreases, and thereby the crystallinity is improved. The high pressure condition is beneficial for suppression of the ion bombardment. The deposition rate and the crystalline volume fraction as a function of the RF power are shown in Figs. 12(a) and (b) for RF and VHF plasmas, respectively. In the low RF power region, the structure is amorphous, and at the onset power of the saturation of the deposition rate the crystalline phase starts to appear. This is due to the increase of the atomic hydrogen density under the depletion condition of silane. In the high power region, the deposition rate decreases due to the etching by atomic hydrogen. The etching is more pronounced for the VHF plasma, suggesting a higher atomic hydrogen density in the VHF plasma. The preferential orientation along [1 1 0] direction is more pronounced in the sample prepared by VHF plasma [33]. This is also due to the lower ion bombardment of the VHF plasma.

0 14

Fig. 11. Pressure dependence of deposition rate and crystalline volume fraction. For (a) RF PECVD and (b) VHF PECVD.

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2

100

Deposition Rate (nm/s)

60 1 40 H2/SiH4 = 5/95 13.56 MHz P = 4 Torr

0.5

20

0 0

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100

150

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80 1 60

40 0.5 H2/SiH4 = 90/10 60 MHz P = 2 Torr

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0 0

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Volume fraction (%)

Deposition rate (nm/s)

0 300

Power (W)

(a)

(b)

Volume Fraction (%)

80

1.5

150

0 200

Power (W)

Fig. 12. The deposition rate and the crystalline volume fraction as a function of the RF power under highpressure depletion conditions using RF plasma (a) and VHF plasma (b).

The Raman spectra as a function of the deposition rate deteriorate significantly at a deposition rate of 5 nm/s even using the high-pressure depletion method (see Fig. 13). This suggests the presence of a disrupting factor in the very high deposition rate region. We have considered three possible causes: ion bombardment, higher silane, and the intrinsic upper limit. The intrinsic limit of the growth rate has been observed in molecular beam epitaxy at low temperatures [34]. However, it turns out that 5 nm/s is not the intrinsic limit. A triode method using a mesh electrode between the cathode and anode electrodes has been developed to reduce the ion bombardment [35]. As shown in Fig. 14, the deposition rate is reduced to less than half by the mesh. With increasing the RF power, however, the deposition rate abruptly increases and reaches 90% of the value for the diode case, accompanied by

Raman Intensity (arb. units)

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5.8 (mesh)

5.0 (HPD)

3.8 (HPD) 2.2 (HPD) 0.006 400

450 500 550 Raman shift (cm-1)

600

Fig. 13. Raman spectra of the film deposited under various methods and deposition rates. The numbers give the deposition rate in nm/s and the deposition method is designated in the parenthesis. The bottom trace is for the film prepared using the conventional method.

a lot of migrating bright spots on the mesh. The film structure is crystalline under such a condition, while only the amorphous phase is obtained without using the mesh. The bright spots involve intense Sin, Ha and Hb emission lines, while no increase in the SiHn intensity was observed, arising from the locally high-density plasma where the decomposition rate significantly increases. Since the spots occur closer to the substrate than the cathode, the deposition rate effectively increases. Another important issue is defects probably in the grain boundary region. For solar cell applications, the recombination of photo carrier is a crucial factor and the recombination is mainly dominated by mid-gap states arising from dangling bonds. The defect density in the film measured by ESR monotonically increases with the deposition rate as shown in Fig. 15. A high efficiency solar cell with ZB10% contains N¼ 2  1016 cm3 or less. As compared to the conventional method, the highpressure depletion method can reduce N by about one order of magnitude probably due to the lower ion bombardment. A further decrease in N is obtained using the triode method. N¼ 2:6  1016 cm3 ; which is nearly one order of magnitude lower at a deposition rate of 5–6 nm/s, as compared to the high-pressure depletion method. Photo- and dark-conductivities are sp ¼ 105 and sd ¼ 107 S/cm, respectively. Usually, the solar cell efficiency is a decreasing function of the deposition rate similarly to the defect density. The high-pressure depletion method is expected to be

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SiH4/H2=30/270 7

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100 without mesh

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80 4 60 3 with mesh

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2 with mesh

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Crystalline volume fraction (%)

Deposition rate (nm/s)

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0 300

250

Power (W) Fig. 14. The deposition rate and the crystalline volume fraction as a function of the RF power under highpressure depletion conditions using VHF plasma with the mesh and without mesh. The film structure is shown only the case with the mesh because the film structure without using the mesh is always amorphous.

1018 (A)

Defect density (cm-3)

(B)

1017

(c) 1016

1015 0

1

2

3

4

5

6

Deposition rate (nm/s) Fig. 15. Defect density measured by ESR as a function of deposition rate using three different methods; (A) conventional method, (B) high pressure depletion method, and (C) triode method.

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30

Current density (mA/cm2)

25 20 15

Area: 0.25cm2 Jsc=23 mA/cm2 Voc=0.514 V FF=0.69 Eff.=8.1 %

10 5 0

0

0.1

0.2 0.3 0.4 Voltage (V)

0.5

0.6

Fig. 16. The I–V characteristics of microcrystalline silicon solar cells prepared at (a) 0.15 nm/s and (b) 1.2 nm/s. The thickness of the i-layer is 2.5 and 2 mm, respectively.

effective to maintain the efficiency for higher deposition rate. The efficiency goes slightly down to 8.1% at 1.2 nm/s as shown in Fig. 16. This number is the highest among those ever reported [36,37]. This result is attributed to the improvement of the film quality. From a material point of view, the structural and electrical properties have sufficient quality for device application by using the high-pressure depletion method. For actual device application, however, we have found a couple of problems; one is the interface damage and the other is the amorphous incubation layer formed at the beginning of the deposition. For a deposition rate around 5 nm/s, a thick amorphous incubation layer is formed on the doped layer at the beginning of the deposition of the i-layer. In order to avoid the formation of this incubation layer, we have attempted to use a two-step growth where a low deposition rate is employed during the initial stage followed by a high-rate deposition. Another advantage effecting other than the film properties is the degree of powder formation. The powder formation involves a higher order reaction through the SiH2 insertion reaction into SiH4. Under the depletion condition, the decrease in SiH4 density may result in a higher silane formation.

6. Effect of surface morphology As the deposition rate is important for a high through-put process, another approach is to enhance the optical confinement in the long wavelength region to

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increase the ‘‘optical’’ thickness of microcrystalline silicon. For p–i–n type amorphous silicon solar cells, a textured SnO2 has been developed. For microcrystalline silicon solar cells, on the other hand, the different wavelength region where the enhancement is required, the absorber thickness, and the texture formed spontaneously during Si growth influence the optimized substrate texture. When two different kinds of transparent conducting oxides (TCO) were compared, one Asahi-U and the other chemically etched ZnO, it was found that a shallower texture is better for mc-Si:H solar cells. The Asahi-U substrate was covered by 50 nm-thick ZnO layer to avoid the reduction of SnO2 by atomic hydrogen. The ZnO layer was deposited using DC-sputtering of a Ga-doped ZnO (ZnO:Ga) target. The other ZnO substrates were textured by the wet etching process using 0.5% HCl water solution, and the surface morphology was varied by etching time [38]. The surface morphology of the substrates was measured by atomic force microscopy (AFM). We obtained the root-mean-squares (RMS) height (h) and an average width of the roughness (l) by Fourier analysis of the surface roughness. Here, we defined an average slope of surface texture tan y by tan y ¼

2h : l=2

ð1Þ

The average slope, tan y , in ZnO substrate deposited on glass, increases with increasing HCl etching time. Fig. 17 shows the solar cell characteristics with a 2.5 mm i-layer plotted as a function of tan y: Jsc increases and then saturates with increasing tan y: This result indicates that efficient light trapping occurs for tan y greater than 0.08. A similar trend has been reported for a-Si:H solar cells deposited on textured substrates [39,40]. It is noteworthy that tan y of the Asahi-U substrate is much higher than the onset of this saturation. On the other hand, Voc monotonically decreases with increasing tan y and FF shows a similar behavior. Because of the balance between Jsc and Voc ; the conversion efficiency shows a maximum at around tan y ¼ 0:08; Z ¼ 9:4% (Voc ¼ 0:526 V, Jsc ¼ 25:3 mA/cm2, FF¼0.710) was obtained as shown in Fig. 16 [41]. The reduction in Voc and FF on steeper substrates can be explained by the difference in microstructures. The values of tan y calculated by Eq. (1) for Figs. 18(a), (b), and (c) are 0.005, 0.084, and 0.3, respectively. The crystalline growth direction is perpendicular to the local surface of the textured substrate as indicated by an arrow in the figure. In the solar cell fabricated on the Asahi-U substrate (Fig. 18(c)), the columnar growth is hindered by collision of the columns, and a large number of grain boundaries can be seen. On the smoother substrates, on the other hand, the collision of the columns is suppressed and the density of grain boundaries is reduced, as shown in Figs. 18(a) and (b). The reduction in Voc and FF with increasing the average slopes of the substrate texture is ascribed to the increased carrier recombination at grain boundary defects. Thus, in mc-Si:H solar cells, optimization of the textured substrate should be performed not only by optical confinement (Jsc ) but also by the control of crystal growth mode (Voc ; FF). It has been pointed out that, when a few tens of nano-meter-size roughness exists at the TCO/p interface, the interface layer can be considered having the middle of

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0.2 Tan


0.3

Efficiency (%)

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8 0.4

Fig. 17. The open-circuit voltage and short-circuit current for various average slope, tan Y (see text) of TCO substrates.

Fig. 18. Cross-sectional TEM images for TCO coated glass substrate with different average slope of (a) 0.005, (b) 0.084, and (c) 0.3, respectively.

refractive index between TCO and p-layer by assuming the effective medium approximation (EMA) [42]. An additional nano-meter-size roughness was examined using island growth at higher temperatures. The lateral size was 130–150 nm and the vertical size was 15–35 nm, respectively. Fig. 19 shows the both quantum efficiency and reflectance of mc-Si:H solar cells fabricated on substrate C (without nanotexturing) and D (with nano-texturing). In the case of the cell fabricated on substrate

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80 Substrate C Substrate D

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70 60 50

0.6

40 0.4

30 20

0.2

Reflectance (%)

Quantum efficiency

1

10 0 400

500

0 600 700 800 900 1000 1100 Wavelength (nm)

Fig. 19. Quantum efficiency and reflectance of mc-Si:H solar cells fabricated on substrate C (without nanotexturing) and D (with nano-texturing).

D; the quantum efficiency in the range between 300 and 700 nm is further improved by the suppression of optical reflectance, compared with that of substrate C: Jsc also increases from 23 to 24.6 mA/cm2 using substrate D; while Voc and FF remain constant. As a result, the efficiency improves from 8.1% to 8.4% (Voc ¼ 0:507 V, Jsc ¼ 24:6 mA/cm2, FF¼0.677). This result suggests that the anti-reflection effect is caused by a few tens of nano-meter-size roughness at the TCO/p interface [43].

7. Interface effects upon solar cell performance mc-Si:H solar cells have two different structures, superstrate and substrate type, where the order of the deposited layer is p–i–n and n–i–p, respectively. The superstrate type mc-Si:H solar cells with Z ¼ 9% can be obtained at temperatures around 1401C [43], while the substrate type cells allow much higher process temperature for obtaining better efficiency [44,45]. The crystallite size increases monotonically up to about 4001C, whereas the defect density shows minimum at around 2501C similarly to a-Si:H [46]. For better carrier collection efficiency, the carrier recombination rates should be lowered by means of an increase of the diffusion length and a reduction of defect density in the i-layer. On the other hand, thermal damage at the p/i or n/i interfaces is crucial for device performance. The presence of impurities is also a crucial factor, and as mentioned in the previous section, oxygen donors can be passivated by low temperature deposition around 140–1801C [27]. Material properties such as defect density and grain size, however, have a higher optimum temperature. The optimum temperature intrinsic to microcrystalline silicon has been studied using highly purified materials prepared by ultrahigh vacuum (UHV)-PECVD [47]. The oxygen concentration in undoped mc-Si:H films measured by SIMS was typically 1  1018 cm3. The carrier

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concentration of i-layer films was measured by an AC-Hall measurement system and was typically 1  1012 cm3. These two values were independent of Ts in the range between 1001C and 2501C. The defect density measured by ESR shows the minimum value of 2  1015 cm3 at Ts E 2501C. Phosphorus and boron concentrations at the interface between undoped layer and doped layers were measured by SIMS. Both n/i and p/i interface properties were studied by means of photoluminescence (PL) measurements in the near infrared region at 10 K. Fig. 20 shows the dependence of solar cell efficiency on the deposition temperature of the i-layer for superstrate configuration with the i-layer thickness of 1 mm and for substrate configuration with a thickness of 2.5 mm. The difference in the efficiency between the substrate and superstrate type cells is ascribed to the i-layer thickness. With elevating Ts of the i-layer, the efficiency of substrate type solar cells is almost constant in this temperature range, while the efficiency of superstrate type solar cells decreases at 2501C. Voc ; Jsc ; and FF show similar behaviors. Carrier density in the i-layer was estimated to be as low as 1  1012 cm3 and was independent of Ts as measured by the Hall effect. The defect density of i-layer films shows the minimum value of 2  1015 cm3 at 250 1C. Therefore, the reduction in efficiency of superstrate type solar cells at 2501C is not ascribed to the i-layer film properties, but to another possibility such as the interface damage between i-layer and the underlying doped layer. The wavelength dependence of the quantum efficiency for various Ts is shown in Figs. 21(a) and (b) for substrate and superstrate type cells, respectively. With elevating Ts of the i-layer, the quantum efficiency in the range between 300 and 700 nm of superstrate cells decreases at 2501C, while those for substrate cells remain unchanged. Since the short wavelength light is almost absorbed near the p/i interface, the reduction of quantum efficiency is ascribed to the poor carrier collection at the p/i interface. The carrier lifetime is estimated by photoluminescence from p/i and n/i structures. The sample structures consist of 200 nm of intrinsic mc-Si and 25 nm of doped mc-Si layers on a synthetic quartz glass substrate. The doped layers were deposited at

Superstrate type Substrate type

Efficiency (%)

9 8 7 6 5 4

100

150 200 250 300 i-layer deposition temperature (°C)

Fig. 20. The dependence of solar cell efficiency on the deposition temperature of the i-layer for superstrate configuration with the i-layer thickness of 1 mm and for substrate configuration with a thickness of 2.5 mm.

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Quantum efficiency

0.8

(a)

0 1

Quantum efficiency

1

0.8

(b)

150°C 180°C 250°C

0.6 0.4 0.2

150°C 180°C 250°C

0.6 0.4 0.2 0 300

400

500

600 700 800 Wavelength (nm)

900 1000 1100

Fig. 21. The quantum efficiency as a function of wavelength for various Ts for: (a) substrate, and (b) superstrate type cells.

1401C and the i-layer was deposited at various temperatures. Fig. 22 shows the intensity of PL arising from the inter-band transition in mc-Si:H plotted for various i-layer Ts : Solid circles and open circles designate the result of n/i samples and that of p/i samples, respectively. The PL intensity for the n/i samples is not sensitive to the i-layer Ts ; while the PL intensity of the p/i sample is much weaker for temperatures higher than 2501C, which is ascribed to the formation of defects acting as nonradiate recombination centers in the near p/i-interface. Fig. 23(a) and (b) shows the depth profile for dopant concentration near p/i and n/i interface for different Ts of the i-layer. The doped layers are prepared at a fixed temperature of 1401C. As shown in Fig. 23(a), the profile of boron has wider spatial extent due to diffusion into the i-layer even at 2501C, which cannot be explained in terms of thermal diffusion of boron itself. In fact, the post-annealing of the 1501C sample at 2501C causes no appreciable difference. This implies that the boron diffusion is facilitated only during the deposition. The profile of phosphorous, on the other hand, is nearly independent of the i-layer deposition temperature as shown in Fig. 23(b). Thus, the enhanced diffusion is observed only for boron and only during the deposition at a higher temperature of 2501C. The enhanced impurity diffusion in amorphous silicon has been reported and has been ascribed to the interaction with hydrogen because of the absence of the

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Intensity (arb. units)

10

1

0.1 n/i 0.01

0.001 140

p/i

160

180

200

220

240

260

i-layer deposition temperature (°C ) Fig. 22. The intensity of PL arising from the inter-band transition in mc-Si:H plotted as a function of the i-layer deposition temperature.

enhancement for sample with a low hydrogen content [48]. The diffusion constant of boron in a-Si:H at 2501C is estimated to be about 1017 cm2/s that is tens order of magnitude higher than that in crystalline silicon, whereas it is too small to account for the boron profile in Fig. 23(a). Boron could selectively diffuse in a-Si:H matrix in the mc-Si layer because of the higher diffusion coefficient. The boron in a-Si:H is known to form a charged defect on the basis of the (8-N) rule [49], resulting in the formation of charged defects, which act as a recombination center. Although the mechanism of the enhancement of the boron diffusion during the deposition is still an open question, the Si–Si bond breaking and its accompanying structural relaxation under impinging atomic hydrogen flux during the i-mc-Si layer deposition may be responsible for the enhancement of boron diffusion.

8. Summary In this chapter, we have reviewed recent developments in the deposition technique, material properties, and device fabrication techniques in microcrystalline silicon prepared by RF (and VHF) plasma CVD. The material properties at high deposition rate over 6 nm/s made tremendous progress during the past 5 years. The material properties such as defect density and crystalline volume fraction show no longer any difference from the high quality material prepared at low deposition rate. The solar cell efficiency for high deposition rate over 1 nm also exceeds the highest value for one tenth of the deposition rate. The remaining problem seems to exist at the interface region and further improvement of the film quality. The defect density becomes more important for tandem cells because the bottom cell thickness could become thicker than 5 mm for

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1021 iB concentration (atoms / cm3)

1020

pµc-Si:H

Si wafer

1019 1018 1017 1016 250°C 150°C

1015 (a) 1014 1021 iP concentration (atoms / cm3)

1020

nµc-Si:H

Si wafer

1019 1018 1017 1016 1015

(b) 1014 0.1

250°C 150°C

0.2

0.3 Depth (m)

0.4

0.5

Fig. 23. The depth profile for dopant concentration near: (a) p/i and (b) n/i interfaces for various deposition temperatures of the i-layer. The doped layers are prepared at a fixed temperature of 1401C.

15% efficiency. Otherwise more sophisticated optical confinement should be developed.

Acknowledgements We would like to thank Drs. A. Matsuda, H. Fujiwara, L.H. Guo, M. Fukawa, S. Suzuki, M. Tanda, T. Wada and Y. Nasuno for invaluable collaboration

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and discussion. We are also indebted to Profs. H. Sugai (Nagoya University), Y. Watanabe, M. Shiratani (Kyushu University) for helpful suggestions. This work is partly supported by NEDO.

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