# N-type multicrystalline silicon wafers and rear junction solar cells

S. Martinuzzi, O. Palais, M. Pasquinelli, F. Ferrazza

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S. Martinuzzi, O. Palais, M. Pasquinelli, F. Ferrazza. N-type multicrystalline silicon wafers and rear junction solar cells. European Physical Journal: Applied Physics, 2005, 32 (3), pp.187-192. 10.1051/epjap:2005085 . hal-03350281

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# N-type multicrystalline silicon wafers and rear junction solar cells

S. Martinuzzi<fnref n="1" />,ª, O. Palais<fnref n="1" />, M. Pasquinelli<fnref n="1" />, and F. Ferrazza<fnref n="2" />

Received: 8 September 2005 / Accepted: 29 September 2005

Published online: 30 November 2005 – © EDP Sciences

**Abstract.** N-type silicon presents several advantages compared to p-type material, among them, the most important is the small capture cross sections of metallic impurities, which are neatly smaller. As a consequence lifetime and also diffusion length of minority carriers should be neatly higher in n-type than in p-type, for a given impurity concentration. This is of a paramount interest for multicrystalline silicon wafers, in which the impurity-extended crystallographic defects interaction governs the recombination strength of minority carriers. It is experimentally verified that in 1.2 Ω cm raw wafers lifetimes about 200 µs and diffusion lengths around 220 µm are measured. These values increase strongly after gettering treatments like phosphorus diffusion or Al-Si alloying. Scan maps reveal that extended defects are poorly active, although in regions where the density of dislocations is higher than 10⁶ cm⁻<fnref n="2" />. Abrupt p+n junctions are obtained by Al-Si alloying after annealing between 850 and 900 °C, which could be used for rear junction cells. Such cells can be processed by means of similar processing steps used to make conventional p-type base cells.

**PACS.** 72.20.Jv Charge carriers: generation, recombination, lifetime and trapping

## 1 Introduction

Today the photovoltaic industry produces p-type base solar cells, which the related technology is very mature, however due to the exponential increase of the cell production the feedstock quantity is limited and we have to face a shortage in the years to come. However, n-type silicon is also of interest because approximately 2000 tonnes/year of n type wastes of Czochralski (Cz) grown Si monocrystals are available and a large part of it, eventually after mixing with low doped silicon may be used in production of n-type multicrystalline ingots (mc-Si) by means of directional solidification. N-type material possesses some remarkable advantages. First, there are little or no boron-oxygen complexes and it has been reported that the lifetime of minority carriers does not decrease with the injection level [1]. As a consequence n-type base cells should work more efficiently at low illumination levels compared to p-type based cells. Secondly, and this is of a paramount importance, the minority carrier capture cross sections of metallic impurities, frequently found in processed silicon, are markedly smaller for holes than for electrons [2]. This property of n-type silicon was clearly explained by Macdonald and Geerligs [3] on the basis of the formation of donor trap levels in the gap which are positively charged or neutral, and can be strongly attractive for

electrons or poorly attractive for holes. Interstitial iron which is one of the more harmful contaminant in mc-Si wafers is a good example because the capture cross sections for electrons and for holes, reported by Istratov [4], are $σ_n = 5 \times 10^{-14}$ cm<fnref n="2" /> and $σ_p = 7 \times 10^{-17}$ cm<fnref n="2" />, respectively. Assuming that Schockley-Read-Hall (SRH) recombination occur and that interstitial iron is the dominant impurity, very high lifetime values are expected for holes in n-type silicon. Although the diffusion coefficient of electrons is 3 times higher than that of holes, it is expected that diffusion lengths of minority carriers, in such a simple configuration, are neatly higher in n-type than in p-type silicon (provided the impurity concentration is the same in both types of wafers). The formation of FeB pairs in boron doped silicon, which are less detrimental for electron lifetime [3, 4], reduces but does not cancel the advantage of n-type over p-type silicon. The same tendency is observed for several other metallic impurities [3], except for chromium, for which the capture cross sections are closed in both n and p type silicon [5]. Obviously, there are other impurities and different sources of recombination centers, like fast and slow diffusers, precipitates and crystallographic defects in mc-Si wafers, and the reality is more complex. In addition, the capture cross sections of minority carriers determined by deep level transient spectroscopy are not those occurring in the base of a solar cell in which the electric field is negligible and the temperature

<footnote n="1">¹ UMR TECSEN, Université Paul Cézanne-Aix-Marseille III, 13397 Marseille Cedex 20, France</footnote>

<footnote n="2">² EniTecnologie, via d'Andrea 6, Nettuno, Italy</footnote>

<footnote>ª e-mail: santo.martinuzzi@univ.u-3mrs.fr</footnote>
is higher than 300 K (this point will be developed later in
this paper).

Today, multicrystalline silicon (mc-Si) wafers or ribbons are the dominant material in photovoltaic industry. Such materials are characterized by the presence of a high density of extended crystallographic defects, like grain boundaries, dislocations and twins, which are more or less associated with segregated impurity atoms. These impurities are metallic fast diffusers (mainly iron, chromium and copper), oxygen atoms or oxygen precipitates and it is the interaction between extended defects and impurities that determines the recombination rate of minority carriers in the wafers. The extended defects, by themselves are little involved in recombinations, and in the large grains of mc-Si, dislocations are the most harmful defects when they are decorated by impurities. Since the capture cross sections of the fast diffusers are smaller in *n*-type silicon, it is expected that the consequences of the impurity-defect interaction are strongly reduced, and this reduction will be enhanced after a gettering treatment because impurity concentrations decrease. However, due to the growth conditions of the mc-Si ingots, i.e. directional solidification of Cz crystal wastes in a mould, oxygen atoms are also present in actual concentrations larger than those deduced from infrared spectroscopy. Indeed such a technique detects only interstitial oxygen atoms while precipitated or aggregated atoms are not taken into account. Oxygen precipitation occurs certainly at dislocation cores, at dislocation clusters and at grain boundaries. These precipitates can, in turn, trap metallic impurities and become recombining, as observed in *p*-type silicon. However, their surfaces are expected to have a hole repelling positive charge, and so in *n*-type materials such precipitates have been found less harmful for minority carriers [6].

Recently, experimental results have reported exceptional values of lifetime in *n*-type wafers cut from ingots or from ribbons: Cuevas et al. [1], Schmiga et al. [7] and Libal et al. [8] have found lifetime values of several hundred of µs in raw materials and values in the range of ms in wafers gettered by a weak phosphorus diffusion.

Such interesting properties can be useful for solar cells
and several papers have been published that dealt with the
use of monocrystalline *n*-type silicon with junctions made
by boron or aluminium in-diffusion. The Ebarra group
has intensively investigated *n*-type monocrystalline silicon
ribbon solar cells in which a rear junction was made by
means of aluminium-silicon alloying [9, 10]. This technique
is very attractive because it is well known that in-diffusion
of aluminium in silicon occurs at relatively low tempera-
tures after the formation at 570 °C of the Al-Si eutectic
alloy, which behaves like a metal. Indeed, during the cool-
ing down of the structure, the alloy retracts while a few
micrometers of silicon remain saturated with aluminium.
A *p*⁺-type layer is formed in the silicon just above the
alloy, and if the alloy is annealed at T > 700 °C, the dop-
ing level reaches 10¹⁹ cm⁻³, i.e. the value of the solubility
limit of Al in Si at such temperatures. As a consequence
an acceptable *p*⁺*n* junction could be obtained more eas-

ily than with boron diffusion, possibly with the additional
benefit of a gettering effect.

The aim of this paper is to show, by means of scan
maps of minority carrier diffusion lengths and lifetimes,
that *n*-type mc-Si is an excellent material, because the re-
combination strength of dissolved impurities and extended
crystallographic defects is low, neatly smaller than in *p*-
type mc-Si, and that the Al-Si alloy formation gives rise
to acceptable *p*-*n* rear junction solar cells (RJC).

2 Experimental

*N*-type 350 µm thick mc-Si wafers were cut from ingots grown at EniTecnologie in a directional solidification furnace, following the usual production conditions for *p*-type ingots. The silicon feedstock was composed of top and tail parts of *n*-type Cz ingots. The resistivity of the investigated wafers measured by the four point probe technique was 1.2 Ω cm, and the *n*-type conductivity was verified using the hot probe test and also by means of a contactless tester (Semilab P/N tester).

To obtain an evaluation of the ratio σ<sub>p</sub>/σ<sub>n</sub> in the base
of a solar cell and to compare its value with that given
by DLTS, we have contaminated voluntarily *p* and *n*-type
wafers (~1 Ω cm) by means of Fe, Co and Au ion implan-
tation at an energy of 100 keV and a dose of 10<sup>13</sup> cm<sup>-2</sup>.
Then the samples are annealed in argon at 850 °C for 2 h,
bulk lifetimes τ<sub>n</sub> and τ<sub>p</sub> are measured and the ratio σ<sub>p</sub>/σ<sub>n</sub>
is evaluated assuming that SRH statistics occurs.

Lifetime and diffusion length (L) of minority carriers
were determined by means of the microwave contactless
phase shift (μWPS) and surface potential voltage tech-
niques, which gave overall values averaged on a few cm².
Using a focused infrared monochromatic (λ = 960 nm)
light spot, the μWPS and the light bean induced cur-
rent techniques (LBIC) allow scan maps of τ and L to
be obtained with a spatial resolution of 50 μm. Notice
that μWPS technique working at a microwave frequency of
10 GHz and at various modulation frequencies of the near
infrared excitation leads to the simultaneous evaluation of
the bulk lifetime and of the surface recombination veloc-
ity, and does not need surface passivation [11]. The LBIC
technique working at various near infra red wavelengths
leads to the determination of L. All these techniques have
been described in details, previously, in references [12,13].
The collecting structures used for the LBIC technique was
Au-Si(n)-Si(n⁺)-Al. A 8 nm thick semitransparent gold
layer was deposited on the *n*-type front surface, while the
back surface was phosphorus diffused (850 °C–20 min) in
order to create a *n*⁺ layer needed to make an ohmic con-
tact. *N*⁺*np*⁺ rear junctions made by means of aluminium-
silicon alloying were used to make rear junction solar cells
(RJC), which the illuminated front surface was passivated
by an *n*⁺ layer due to phosphorus diffusion. The *p*⁺*n* junc-
tion was obtained after deposition by electron gun evapo-
ration of a 1.2 μm thick layer of aluminium and annealing
at 850 °C for 10 min. The metallisations of the cells re-
sult from the back Al-Si alloy, which is quasi metallic, and
<figure><img src="image_2.png" /><figcaption>Fig. 1. Structure of a rear junction solar cell.</figcaption></figure>

<figure><img src="image_3.png" /><figcaption>Fig. 2. Typical LBIC contrast scan map at $\lambda = 980$ nm of a raw sample with local values of minority carrier diffusion length in $\mu$m.</figcaption></figure>

from a front silver grid deposited by vacuum evaporation. The structure of the cell is described in Figure 1. RJC cells were also prepared using a back side screen printed aluminium layer in order to simulate industrial processing steps used to make $n^+p^+p^+$ conventional cells.

Chemical etching was used to reveal the presence of extended defects in the investigated samples. Grabbing treatments consisted in phosphorus diffusion from $POCl_3$ at 900 °C for 1 h.

## 3 Results

In the raw n-type material the mean values of $\tau_p$, averaged in 25 cm$^2$ wafers, are around 100 $\mu$s. After grabbing they reach, at least, 300 $\mu$s, in agreement with Cuevas et al. [1], Schmiga et al. [7] and Libal et al. [8]. The mean diffusion length values are also very high, as 200 to 300 $\mu$m are measured in raw wafers. Such high values have never been observed in p-type mc-Si raw wafers, or they had been measured only after long grabbing treatments by phosphorus diffusion, or aluminium-silicon alloy grabbing [14].

In phosphorus grabbed samples, the values of $L_p$ are higher than 500 $\mu$m, in agreement with the results of Cuevas et al. [1]. A typical LBIC contrast scan map of

<figure><img src="image_4.png" /><figcaption>Fig. 3. Lifetime scan map of a sample containing dislocation arrays, grain boundaries and subgrain boundaries. Notice that the lower value is 60 $\mu$s in the central area (demarcated by a rectangle) although the extended defect density is very high.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>Fig. 4. Scan map of the LBIC contrast of Figure 3 sample at 980 nm. (contrast is intensified in order to reveal extended defects). Some values of $L_p$ are indicated in $\mu$m.</figcaption></figure>

a raw wafer is given by Figure 2 and some local values of $L_p$ are indicated in $\mu$m. Conversely to p-type materials, we don't have found regions in which there is no improvement of the material and where the diffusion length remains very poor (few tens of $\mu$m).

Figures 3 and 4 show the scan maps of lifetime and of LBIC contrast for a sample in which extended defects are present in high densities, as revealed by the microphotography of Figure 5. In the LBIC map some values of $L_p$ are indicated in $\mu$m.

Notice that the local values of $\tau_p$ are at least ten times higher than those found in a p-type sample containing similar defect densities. Table 1 gives the values of lifetime measured in n and p-type samples, contaminated under the same conditions, and the deduced values of the capture cross section ratio $\sigma_n/\sigma_p$. This ratio is different of that obtained by DLTS [2,3], but it is sufficiently large for Fe, for FeB and for Co in order to confirm the advantage of n-type silicon and to check that the values given by DLTS are acceptable. Both maps and the
<figure><img src="image_6.png" /><figcaption>**Fig. 5.** Optical microphotography of the central part of Figure 3 sample after chemical etching. The lower values, 60 µs for $τ_p$ and 180 µm for $L_p$, are found in regions where the dislocation density is higher than $10^6$ cm$^{-2}$.</figcaption></figure>

<table><caption>**Table 1.** Values of the ratio of capture cross sections for minority carriers obtained from DLTS and from lifetime measurements at room temperature in *n* and *p* type mc-Si samples contaminated under the same conditions. In *p*-type $τ_n$ is measured before and after dissociation of FeB pairs.</caption><thead><tr><th>Impurity</th><th>σ (cm<sup>2</sup>)</th><th>τ (μs)</th><th>σ<sub>n</sub>/σ<sub>p</sub> (DLTS)</th><th>σ<sub>n</sub>/σ<sub>p</sub> (τ)</th></tr></thead><tbody><tr><td rowspan="2">Fe<sub>i</sub></td><td>σ<sub>n</sub> = 5 × 10<sup>-14</sup></td><td>τ<sub>p</sub> = 300</td><td rowspan="2">700</td><td rowspan="2">300</td></tr><tr><td>σ<sub>p</sub> = 7 × 10<sup>-17</sup></td><td>τ<sub>n</sub> = 1</td></tr><tr><td>FeB</td><td>σ<sub>n</sub> = 2 × 10<sup>-15</sup></td><td>τ<sub>n</sub> = 20</td><td>28</td><td>15</td></tr><tr><td rowspan="2">Au<sub>s</sub></td><td>σ<sub>n</sub> = 2.2 × 10<sup>-15</sup></td><td>τ<sub>p</sub> = 25</td><td rowspan="2">18</td><td rowspan="2">4</td></tr><tr><td>σ<sub>p</sub> = 1.4 × 10<sup>-16</sup></td><td>τ<sub>n</sub> = 6</td></tr></tbody></table>

microphotography reveal the same features of recombining defects and their comparison shows that the high values of $L_p$ are due to the high values of $τ_p$. In a region of the sample of Figure 3, demarcated by a rectangle and in which the dislocation density is about $10^7$ cm$^{-2}$, it is found that $L_p \sim 120$ µs and $τ_p \sim 60$ µs. Such values confirm that the recombination strength of the extended defects is very low.

With this excellent material we can realize easily junctions for solar cells by Al-Si alloying, as said in the introduction. Because the alloy is electrically conductive but not transparent to sunlight, rear junction cells (RJC) must be used, which work efficiently provided that the minority carrier diffusion length is higher than the substrate thickness and that the front surface is passivated (by means of a slight phosphorus diffusion).

Computed variations of the conversion efficiency as a function of the base thickness for various $L_p$ values and for a passivated illuminated surface (recombination velocity $S = 0$ cm/s) is given by Figure 6. Using thin *n*-type multicrystalline silicon wafers, which the thickness is smaller than 120 µm, in which $L_p$ reaches 2 to 3 times the wafer thickness, the conversion efficiency of the RJC can reach 19%. As the passivation level of the illuminated surface is critical, we have also computed the variation of the conversion efficiency as a function of $L_p$ for various values of

N$^{+}$NP$^{+}$ Rear Junction Cell

<figure><img src="image_7.png" /><figcaption>**Fig. 6.** Conversion efficiency versus base thickness of a RJC for various diffusion length $L_p$ and surface recombination velocity $S = 0$ cm/s.</figcaption></figure>

N$^{+}$NP$^{+}$ Rear Junction Cell

<figure><img src="image_8.png" /><figcaption>**Fig. 7.** Conversion efficiency of a RJC versus $L_p$ for various values of $S$.</figcaption></figure>

$S$ and a wafer thickness of 200 µm. Figure 7 shows that the conversion efficiency does not change significantly when $S$ is below $10^3$ cm/s. We have verified that values of $S$ lower than 100 cm/s are obtained after a front surface phosphorus diffusion at 850 °C for 20 min and that acceptable $p^n$ junctions are obtained after annealing of the Al-Si structures at temperatures in the range 850 to 900 °C. The diffusion potential is around 0.8 eV and recombination and diffusion density currents are about $10^{-6}$ A/cm$^2$ and $10^{-8}$ A/cm$^2$, respectively, in agreement with the results of Ebarra group [9,10].

The photovoltaic parameters of 4 cm$^2$ RJC cells made with a 300 µm thick raw wafer and an evaporated aluminum layer are: $V_{oc} \sim 560$ mV and $J_{sc} \sim 16$ mA/cm$^2$ without antireflection coating. Notice that for such a solar cell PC1D predicts a photocurrent density around 17 mA/cm$^2$, which is in good agreement with the experimental value.

Using a 200 µm thick wafer (chemically thinned) the photocurrent reaches 24 mA/cm$^2$. With a screen printed Al layer, a 300 µm thick wafer and a 22 cm$^2$ RJC we measure $V_{oc} \sim 590$ mV and $J_{sc} \sim 16$ mA/cm$^2$. As shown by Figure 8, the LBIC technique works well at $\lambda = 960$ nm when the cell is illuminated by the phosphorus diffused
<figure><img src="image_9.png" /><figcaption>Fig. 8. LBIC contrast scan map of a RJC at $\lambda = 960$ nm.</figcaption></figure>

surface, although the thickness of the wafers is too large. The images of the grain boundaries are enlarged because the minority carriers must cross the entire wafer thickness before to be collected by the rear junction, but the grains appear very homogeneous, confirming that intra grain defects (dislocations and impurities) are poorly recombining.

## 4 Discussion

The preceding results confirm the high quality of *n*-type raw mc-Si wafers which exhibit very large values of lifetime and diffusion length of minority carriers. The high values of $L_p$ are directly connected to the high values of $\tau_p$ since the hole diffusion coefficient in the grains, computed from the expression $L = (D\tau)^{1/2}$ is about 10 cm²/V s, i.e. close to the conventional value in silicon. There is no doubt that it is due to the small capture cross section of metallic impurities which reduces the recombination strength of dissolved, precipitated and segregated impurities. As a consequence crystallographic defects, especially dislocations and subgrain boundaries, are poorly active.

The capture cross section ratio evaluated in *p* and *n* type mc-Si wafers from lifetime is smaller than that deduced from DLTS measurements carried out in monocrystals, this can be explained by the influence of native impurities present in mc-Si, which is not negligible compared to that introduced by ion implantation. As an example, interstitial chromium (frequently found in mc-Si wafers) reduces the advantage of *n*-type material because $\sigma_n \sim \sigma_p$ (2). Fe and Cr tend to form pairs with boron in *p*-type material which the capture cross sections are different from those of interstitial atoms. Moreover these pairs can be partly dissociated by sunlight irradiation leading to a very complex situation. Nevertheless the lifetime values measured in both type of mc-Si lead to a neat advantage for *n*-type material which could be very useful in solar grade material made from purified metallurgical silicon.

The high sensitivity of *n*-type wafers to gettering treatments, like phosphorus diffusion, is not due to particular properties related to the type of conduction or to the presence of phosphorus atoms as dopants. Like in *p*-type mc-Si

wafers the gettering treatments reduces the density of fast diffusers. Even if the concentrations of such impurities are of the same order in *n*-type as in *p*-type silicon after a given gettering treatment, their recombination strength is weaker in *n*-type, and the influence of impurity-defect interaction on the recombination strength of extended defects is smaller. As a consequence $\tau_p$ and $L_p$ achieve very high values in the raw material [1].

The results obtained for our solar cells are limited by the too large thickness of the raw wafers which reduces the photocurrent. $L_p$ is not sufficiently large and gettering treatments are necessary when the wafer thickness exceeds 200 µm. This assumption is confirmed by the agreement obtained between the measured photocurrent density and the value computed by PC1D. Moreover the results show that technical processes used to make solar cells with *p*-type wafers (phosphorus diffusion; screen printing of aluminium and aluminium-silicon alloying) work with *n*-type material, and that is an advantage of RJC (notice that for *n*-type base solar cells phosphorus diffusion passivates the front surface and contributes to obtain a good ohmic front grid, while backside Al-Si alloying forms the $p+n$ junction).

## 5 Conclusion

The results presented in this paper confirm that *n*-type multicrystalline silicon is a high quality material which the lifetime and diffusion length of minority carriers reach very high values: in raw wafers the mean values, averaged on a 25 cm² area, are in the range 80 to 120 µs and 200 to 250 µm, respectively. These values are still higher in gettered wafers, after phosphorus diffusion or aluminium-silicon alloying. This results essentially from the low values of the recombination center capture cross sections due to metallic impurities. The scan maps of diffusion length and of lifetime of minority carriers show that the effects of the interactions between impurities and extended defects are strongly reduced in *n*-type multicrystalline silicon, compared to *p*-type, and consequently the recombination strength of grain boundaries and of dislocations, although their concentrations could be very high in some sample regions, is neatly lower.

Such material could be used to make solar cells with a rear junction made through aluminium doping after aluminium-silicon alloying.

The authors would like to thank the French Agency for Environment and Energy Management as well as CNRS-ECODEV, France, for their support.

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