Proceedings of 8th Thin Film Materials & Devices Meeting
November 4-5, 2011, Kyoto, Japan
http://www.tfmd.jp/

Surface Passivation of Crystalline Silicon by Micro Crystalline Silicon
Deposition Followed by High-Pressure H2O Vapor Heat Treatment
Tomokazu Nagao1, Toshiyuki Sameshima1,
Yasunori Andoh2, Asuka Shuku2 and Eiji Takahashi2
1.Tokyo University of Agriculture and Technology, 2-24-16, Nakachou, Koganei, Tokyo
184-8588, Japan
2.NISSIN Electric Co., Ltd, 47, takazechou, umezu, Kyoto, 615-8686, Japan
E-mail : tsamesim@tuat.ac.jp
Keywords: Micro crystalline silicon, H2O vapor heat treatment, Passivation, Minority carrier
lifetime, Microwave absorption system.

Abstract. We report passivation of the crystalline silicon surface by deposition of undoped
microcrystalline silicon (c-Si) followed by high-pressure H2O vapor heat treatment.
25-nm-thick undoped c-Si films were formed to the surface of 500-m-thick n and p-typesilicon substrates using inductively coupled plasma chemical vapor deposition (ICP-CVD) at
150 oC. Samples were heated at 180 and 260oC with 1x106-Pa-H2O vapor for 3h. The
effective minority carrier lifetime eff was measured by 9.35-GHz microwave absorption system
in order to evaluate surface passivation effect. High eff were achieved as 1.8x10-4 and 5.8x10-5s, respectively, by formation the c-Si films on the n and p-type silicon surfaces, respectively.
Moreover, eff were further increased to 1.4x10-3 and 3.4x10-4 s, respectively, by subsequent
high-pressure H2O vapor heat treatment at 260 oC for n and p-type silicon samples.

Introduction
Microcrystalline silicon (c-Si) films can be formed at low temperatures between 100
and 500 oC by plasma-enhanced-chemical-vapor deposition (PE-CVD). Taguchi et al. reported
capability of surface passivation using c-Si films [1]. SiNx and SiOx insulating layer have been
widely applied to surface passivation for crystalline silicon (c-Si) solar cells. Passivation of
metal silicon interface is also important to effectively transporting photo-induced carriers
because direct deposition of metal electrodes to the c-Si surface, markedly increases carrier
recombination velocity and decrease the conversion efficiency of solar cells. SiNx and SiOx
insulating thick films cannot be used. The other hand, c-Si has low band gap energy and highconductivity compared with SiOx or SiNx. Therefore, c-Si films will be suitable for the
conductive passivation between metal electrodes and c-Si substrates. Moreover, c-Si films
have high impurity activation property compared with amorphous silicon. Therefore,
doped-c-Si will be useful for fabrication of high performance crystalline silicon solar cells,such as Hetero-Junction with Intrinsic Type Silicon Solar Cells [2]. However, there has been
still a problem of the plasma damage by PECVD.
In this paper, we report surface passivation of c-Si film deposition followed byhigh-pressure H2O vapor heat treatment. We have reported that high-pressure H2O vapor heat
treatment eliminates plasma damage and the defects in amorphous silicon (a-Si) films and at- 120223055-01 -polycrystalline grain boundaries [3,4]. We report that high-pressure H2O vapor heat treatment
effectively decreases the carrier recombination velocity at the c-Si /c-Si interface.

Experiment
We prepared n and p-type (1 0 0)
oriented CZ-silicon substrates with athickness of 500 m coated with
100-nm-thick thermally grown SiO2
layers (initial samples). The resistivityof n and p-type silicon substrates was
35 and 25 cm, respectively. The SiO2
layers at the top surface were thenremoved using 5%-diluted hydrofluoric
acid. We kept the samples for threedays to form the native oxide layer at
the top surfaces in air atmosphere atroom temperature [5] to prevent
Fig.1 : Schematic of microwave absorption system
epitaxial growth of crystalline silicon
during PECVD. Optical reflectivity
measurement revealed that the average thickness of native oxide layers was thin about 1 nm,
which was enough tunneling quantum electrical conduction. Undoped c-Si films with athickness of 25 and 3 nm were deposited at the top surfaces at 150oC by ICP-CVD at 0.5 Pawith mixed gases with SiH4 at a flow rate of 50 Sccm and H2 at 25 Sccm [6]. For comparison,
we formed 25-nm-thick c-Si films immediately after 5%-diluted-hydrofluoric-acid treatment.
The Samples were subsequently heated in 1x106-Pa-H2O vapor at 180 and 260oC for 3h.
eff was measured by 9.35-GHz microwave absorption system, as shown in Fig. 1 [7,8].
The sample substrates were put in a gap of wave-guide tubes. Microwave transmittance signalsof the samples were measured when continuous wave or periodically pulsed light with 620
nm-wave length was illuminated to the top surfaces of samples at 1.5 mW/cm2. Moreover, thesignals were coincidentally integrated with time during light ON to reduce the noise level. eff
was estimated using the carrier diffusion and annihilation theory. The surface recombination
velocity at the top surfaces Stop was estimated from eff using equation (1) based on the carrier
diffusion and annihilation theory with the minority carrier diffusion coefficient D and thethickness of substrate d. We assumed that bulk lifetime b in the c-Si substrates was very long at
1 s and the recombination velocity at the rear surface Srear for n and p-type samples kept thesame initial value of 17 and 100 cm/s obtained from initial eff of 1.5x10-4 and 2.5x10-4 s,
respectively. The optical penetration depth for c-Si at 620-nm-wave length was about 2.2 m,which was much lower than the substrate thickness. Therefore, we assumed that carrier
generation was limited in the top surface region.


 D

D
d  Dd 
 Srear   Srear  exp  
 1  exp  





 b 
D b     b
D b  
 b



 eff   b

 D
 D
  D
 D

2d 
 Srear   Stop   
 Stop   Srear  exp  D b  b b  b b- 120223055-02 -(1)

Results and Discussion
Figure 2 shows Raman scattering
spectrum for c-Si films formed on glass
substrates. This spectral curve was analyzed
by three Gaussian curves with peakwavenumber of 470, 500, 520 cm-1,
associated with bonding vibrations ofamorphous, nano-crystalline and crystalline
silicon, respectively. The best fitting resulted
in a crystalline volume ratio of 0.5 from theratio area of nano-crystalline and crystalline
Fig. 2 : Raman scattering spectra for c-Si
components to the total area.
Experimental eff are summarized inFig. 3 (a) and (b) in the cases of 25 and3-nm-thick c-Si films, respectively, for n
and p-type silicon samples.
Stop wereanalyzed from experimental eff. They aresummarized in Fig. 4 (a) and (b) in the casesof 25 and 3-nm-thick c-Si films,
respectively, for n and p-type silicon samples.
The samples with bare top surfaces kept for
three
days
after
after
5%-diluted-hydrofluoric -acid treatment had
low eff of 2.0x10-5 and 5x10-6 s for n andp-type samples. Stop were therefore high of2500 and 7500 cm/s, respectively. On theother hand, 25-nm-thick c-Si deposition
increased eff to 1.8x10-4 and 5.8x10-5 s for n
and p-type samples, respectively, as shownin Fig.3 (a). Stop decreased to 260 and 730
cm/s, respectively, as shown in Fig.4 (a). We
interpret native oxide does not have stable
SiO2 bonding network and that nativeoxide/silicon interface is also unstableaccompanied with many defects. Moreover,native oxide was spontaneously andinhomogeneously formed on the siliconsurface. We believe that there must be someclucks between native oxide domains, which Fig.3 : Minority carrier lifetime eff for ncause carrier recombination sites. We believe (circles) and p-type (triangles) samples withthat silicon deposition it does not work 25(a) and 3(b) nm-thick-c-Si films at
perfectly passivation of the silicon surface. as-top-SiO2 removed, as-c -Si deposited,
Silicon deposition was effective to increase and final H2O vapor heat treatment at 180
eff probably because the recombination and 260oC.
velocity was decreased at the region cluck
regions, in which we believed amorphous silicon was formed just at silicon surface. c-Si
grains grew on the a-Si incubation region. Subsequent 1.0x106-Pa H2O vapor heat treatment at
260oC for 3h markedly increased eff for n and p-type samples to 1.4x10-3 and 3.4x10-4 s,
respectively. Stop decreased to 19 and 51 cm/s, respectively. Those results indicate that- 120223055-03 -high-pressure H2O vapor heat treatment
effectively decreased the density ofrecombination sites at a-Si /c-Si [9] and nativeoxide/c-Si interfaces probably because theirinterfaces became stable. eff for n and p-typesamples were finally comparable to values of
the initial samples with the both surface
coated with 100-nm thermally grown SiO2
layers. Moreover, 180oC H2O vapor heat
treatment also increased eff for n and p-typesamples with 25-nm-thick c-Si to 1.3x10-3
and 3.0x10-4 s, respectively. The results oflong eff obtained by 180oC H2O vapor heat
treatment shows that the combination c-Si
film formation with high pressure H2O vapor
heat treatment is effective for surface
passivation at low process temperature.
On the other hand, eff was increased
to 4.7x10-5 and 1.5x10-5 s by 3-nm-thick c-Si
deposition for n and p-type samples, as shownin Fig. 3(b). However, high pressure H2Ovapor heat treatment was not effective tofurther increase eff. They were 5.6x10-5 and1.7x10-5 s for n and p-type samples,
respectively when 1.0x106-Pa H2O vapor heat
treatment at 260oC for 3h was carried out. We
interpret that the surface 3-nm-thick c-Si
was not enough to completely cover the c-Si
surfaces.
Moreover, eff of the n-type siliconsample coated with 25-nm-thick c-Si
immediately after 5%-diluted-hydrofluoric-acid treatment had a low value about 1.3x10-5
s. High-pressure H2O vapor heat treatment at
260 oC did not increase eff. This result
indicated that Pre-formed thin oxide layers
before c-Si deposition was important toformed high quality c-Si/c-Si interfaces. We
probably thought that the native oxide layer
protected c-Si surface from plasma damage.
Figure 5 shows changes in eff as afunction of elapse time for n and p-typesamples coated with 25-nm-thick c-Si films
treated with 1.0x106-Pa H2O vapor heat
treatment at 260oC for 3h, were kept in air
atmosphere at room temperature. eff kept highvalues of 1.1x10-3 and 2.5x10-4 s for n andp-type samples, respectively, in 150 days after
high-pressure H2O vapor heat treatment. Those

Fig.4 : Surface recombination velocity Stop
analyzed from experimental eff for n andp-type samples with 25(a) and 3(b)nm-thick-c-Si films at as-top-SiO2
removed, as-c -Si deposited, and final H2Ovapor heat treatment at 180 and 260oC.

Fig.5 : The change in minority carrier
lifetime eff for n (circles) and p-type(triangles) as H2O vapor heat treatment at
260 oC samples with elapse time.

- 120223055-04 -results of Figs. 3, 4 and 5 show that the present method of c-Si film formation followed byhigh-pressure H2O vapor heat treatment gives effective and stable passivation of the siliconsurfaces.

Summary
Passivation of crystalline silicon surfaces was investigated by c-Si film formation
using ICP-CVD. Long eff of 1.4x10-3 s and 3.4x10-4 s were achieved for 500-m-thick n andp-type crystalline silicon substrates by 25-nm-thick c-Si deposition followed by 1x106 Pa H2Ovapor heat treatment at 260oC for 3h. Those eff were comparable to the value of 1.5x10-3 and2.5x10-4 s for samples coated with 100-nm-thick thermally grown SiO2 layers. Moreover,
180oC H2O vapor heat treatment also increased eff for n and p-type samples with 25 nm-thickc-Si to 1.3x10-3 and 3.0x10-4 s, respectively. Those results of long eff shows that the
combination c-Si film formation with high-pressure H2O vapor heat treatment is effective forsurface passivation at low process temperature.

Acknowledgments
This work was partially supported by Grant-in-Aid for Science Research C (No.
22560292) from the Ministry of Education, Culture, Sports, Science and Technology of Japan,
and Mikiya Science and Technology foundation.
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