E L Gurevich
email: evgeny.gurevich@isas.de

S Kittel

R Hergenröder

Yu A Astrov
email: yuri.astrov@mail.ioffe.ru

L M Portsel

A N Lodygin

V A Tolmachev

A V Ankudinov

Modification of GaAs surface by low-current Townsend discharge

Keywords: numbers: 52.80.Dy, 68.47.Fg, 52.77.Bn, 68.37.Ps

Influence of stationary spatially-homogeneous Townsend discharge on the (100) surface of semi-insulating GaAs samples is studied. Samples exposed to both electrons and ions in nitrogen discharge at the current density j=60 µA/cm 2 are studied by means of X-ray photoelectron spectroscopy, ellipsometry and atomic force microscopy. It is shown that exposure to low-energy ions (< 1 eV) changes crystal structure of the semiconductor on the depth up to 10-20 nm, although the stoichiometric composition does not change. The exposure to low-energy electrons (< 10 eV) forms an oxide layer, which is 5-10 nm thick. Atomic force microscopy demonstrates that the change in the surface potential of the samples may exceed 100 mV, for the both discharge polarities, while the surface roughness does not increase.

Introduction

Interaction of various materials, such as metals, semiconductors, dielectrics and polymers with gas discharge plasma leads to changes in their behaviour due to the surface modification. This phenomenon is widely used in semiconductor technology for etching, film deposition and modification of surface properties. In gas-discharge devices, however, change in the electron emission properties of cathodes under the action of gas discharge plasma may be undesirable, while it leads to instabilities of discharge parameters in time.

Changes in the semiconductor-electrode properties under interaction with gas discharge plasma is also observed, when studying semiconductorgas discharge (SGD) structures. The structure is a main part of an infrared (IR) image converter, which was invented in the beginning of 70 th [START_REF] Sh | Device for tracking images Russian Authors[END_REF] but similar ideas were discussed even 40 years before [START_REF] Lengnick | Discharge Tube United States Patent Office[END_REF]. The device includes a thin gas discharge gap, one electrode made of a photosensitive in the IR spectral range semiconductor and the second electrode made of a glass plate covered with a transparent conductive film [START_REF] Portsel | [END_REF]. The device is specified by a good sensitivity and high speed of operation [START_REF] Marchenko | Proc. SPIE[END_REF]. However, there may be observed a noticeable drift of its parameters during long operation in the continuous mode. Apparently, it is connected with change in semiconductor surface properties during the long-term action of the discharge plasma on the semiconductor electrode of the device. Such changes are observed if e.g. GaAs semiconductor wafers are used for the electrode fabrication.

The aim of the current work is to study changes in the GaAs surface, which occur after plasma treatment of the semiconductor electrodes at Townsend discharge conditions. The experiments were held in nitrogen atmosphere at low current density. The discharge effects were monitored by measuring optical properties and chemical composition of the surface by means of X-ray photoelectron spectroscopy (XPS) and spectroscopic ellipsometry (SE) before and after the plasma treatment. Besides, roughness of the sample surface and the surface potential were studied by means of atomic force microscopy (AFM). The discharge influence was studied for both polarities of the applied feed voltage, i.e. the semiconductor surface in different experiments was treated by positive and negative (electrons and negative ions) charge carriers.

Considerable changes in the composition of the surface layer of GaAs wafers were observed at a low-current Townsend discharge even for low doses of discharge action. The discharge action was quantified as the deposited charge density, which can be calculated as product of the discharge current density and the discharge time. Oxidation of the electrode surface was found when the semiconductor electrode acted as an anode, i.e. as a result of electron irradiation. Otherwise, when semiconductor electrode served as a cathode (which corresponds to the case of ion exposure), the presence of a thin layer of structurally changed material was revealed. The AFM studies clearly demonstrated that Townsend discharge does not increase the surface roughness even for a stronger plasma treatment, while the surface potential of GaAs changes considerably.

Experimental

The experimental cell is a sandwich-like structure schematically shown in Fig. 1. The gas-discharge gap is between two electrodes, one of which is a SnO 2 -covered glass plate. The other electrode is a semiconductor wafer. The gas-discharge gap is formed by an opening in the dielectric spacer, placed between semiconductor wafer and the glass plate.

The discharge gap width is determined by the spacer thickness, which is 90 µm.

The discharge influence on semiconductor wafers were studied using a semiinsulating (SI) GaAs. Such wafers are used in semiconductor industry as substrates for epitaxy. In all the experiments, the (100) surface of a GaAs sample served as a plane electrode to the gas-discharge. A semi-transparent to visible light electrical contact on a semiconductor plate was formed on the back side of the plate by means of nickel evaporation.

The experimental chamber with the gas discharge cell inside was pumped out and then filled with nitrogen at the residual pressure of 1×10 4 Pa. Nitrogen gas was produced by evaporation of liquid nitrogen with consequent additional purification by means of a liquid-nitrogen trap. The cell was fed with high voltage by a DC power supply. The breakdown voltage of the discharge gap was approximately 300 V.

When amplitude of the feeding voltage applied to electrodes of the device exceeds the breakdown value, discharge current flows through the hybrid SGD structure. The value of the current is controlled by resistance of the semiconductor electrode, which can be adjusted by applying homogeneous illumination from an incandescent lamp. Application of high-ohmic electrode material along with relatively small gas-discharge gap and current density provides formation of stable Townsend discharge. This discharge mode allows spatially homogeneous distribution of the discharge current density over the electrode plane and hence spatially homogeneous irradiation of the studied sample surface. The macroscopic homogeneity of the discharge over the discharge area was monitored by means of a CCD camera.

Electric field accelerates electrons and ions towards the electrode surfaces of a gasdischarge system. Ion interaction with the cathode is especially important, since this process is responsible for generation of secondary electrons, hence for the charge-carriers balance in the discharge. In this paper we study results of the electrode surface exposure with ions and electrons. For that, the polarity applied to the semiconductor electrode was changed in different experiments between negative (in case of its exposure with ions) and positive (used to expose the sample with electrons). Experiments were made at the discharge current density j = 60 µA/cm 2 . To obtain the desired current density, intensity of light illuminating the semiconductor electrode was adjusted. Such a low value of current density guarantees operation of the discharge in the Townsend mode, since it is much lower than the critical current of the Townsend-glow transition [START_REF] Yu | Gas Discharge Physics[END_REF]. The exposure time was in the range of 15-100 minutes, so density of electric charge passed through the system during the exposure (used as a measure of the discharge effect) was in the dose range of Q = 0.05-0.50 C/cm 2 .

The chemical composition was analysed by means of X-ray photoelectron spectroscopy (XPS). This technique identifies chemical composition of studied samples according to energy spectra of photoelectrons. The electrons are emitted from corelevel shells after ionization by soft X-rays. Since the atomic states are influenced by neighbouring atoms, the XPS-technique allows identification of chemical compounds also. The used instrument was an Axis HS from Kratos Analytical, which was equipped with an ion gun for depth profiling of samples. The spectra have been recorded under the following conditions: excitation with non-monochromatic MgKα radiation, the hemispherical analyser was operating in the fixed analyser transmission (FAT) mode with a pass energy of 80 eV under a take-off angle of 0 with respect to the surface normal. Operating in the magnetic mode, the diameter of the measured area was 150 µm.

For the depth profile, a beam of Ar + -ions with an energy of 3 keV was used. Its current was 250 nA and the angle of incidence was 45 with respect to the surface normal. The sputtered area was 2 mm×2 mm. The relative sputtering rate was calibrated with a 100 nm Ta 2 O 5 reference sample from the National Physical Laboratory. Under these conditions, the sputtering rate was 1.0 nm/min referred to Ta 2 O 5 . Etching cycles of 1 minutes were chosen for in-depth profiling, which corresponds to approximately 1 nm per cycle.

For each sample processed in the gas discharge and for the untreated sample, a wide-scan spectrum and narrow scan spectra for O 1s, C 1s, Ga 3d and As 3d were acquired for each analysed depth. The carbon on the surface is due to normal surface contamination. C 1s peak was strongly interfered by the Ga LMM Auger peak. Even though a contamination peak of C 1s could be observed at the surface, a quantification of the C 1s signal under these conditions was not useful. For oxygen, the whole O 1s signal was quantified as one peak.

Comparative narrow scan spectra for arsenic and gallium 3d peaks can be found in Figures 2 and3, respectively; curves corresponding to different depths are shifted vertically for clarity. Numbers at the curves indicate the corresponding sputtering time in seconds; the resulting in-depth step is approximately 2 nm. In case of gallium two components were fitted: GaAs at approximately 19 eV binding energy and Ga 2 O 3 with a constant shift of +1.2 eV. In Fig. 2 one can clearly see that near the surface the peak is located at approximately 20.2 eV, which corresponds to the gallium oxide. However the oxide peak becomes weaker and the GaAs peak at 19 eV stronger, while analysing layers being deeper under the sample surface. A similar behaviour can be observed for arsenic in Fig. 3, where three components are detected and fitted: As bound as GaAs at 41 eV and as oxides As 2 O 3 at 44.5 eV and, in some cases, as As 2 O 5 at 46 eV binding energies. One can see that the oxide dominates only in the surface layer, which is 6-8 nm deep.

Common pre-quantification procedure was applied to the measured XPS spectra: First, the individual peaks were Gaussian-fitted and then integrated after the base line subtraction. Quantification of the spectra has been done with relative sensitivity factors known from literature [START_REF] Briggs | Practical Surface Analysis, Auger and X-ray Photoelectron Spectroscopy[END_REF] corrected for the particular instrument used here. Due to the low binding energy of Ga 3d and As 3d levels, values had to be extrapolated. For this reason, the quantification has to be treated as semi-quantitative. Nevertheless the depth profiles indicate the trend of the chemical changes which are made during the irradiation with electrons or ions as a function of the depth.

For each analysed depth the chemical composition is calculated according to equation ( 1) below. The content C i of an i th component is proportional to the corresponding measured XPS peak area I i divided by the corresponding relative sensitivity factor k i . The result is normalized to the sum of all measured components:

C i = I i /k i j I j /k j × 100% (1)
Optical properties of wafer surface treated with the gas-discharge plasma were investigated by spectroscopic ellipsometry (SE). Ellipsometry measures change in light polarization as a result of reflection from a sample surface at a non-normal incidence. The change in polarization is given by a complex number ρ related to the sample reflection properties as:

ρ = R p R s = tan Ψe i∆ , (2)
where R s and R p are complex reflection coefficients for light waves with the electric field polarized parallel and perpendicular to the plane of incidence, while Ψ and ∆ are the ellipsometric parameters of the polarized light, measured in experiments. In spectroscopic ellipsometry ρ is measured as a function of photon energy hν at a fixed angle of incidence φ 0 . For convenience, ρ-data are sometimes converted to the complex pseudo-dielectric (PD) function, assuming a two-phase (ambient-substrate) model [START_REF] Azzam | Ellipsometry and Polarized Light[END_REF]

ε = ε 1 + i ε 2 = sin 2 φ 0 1 + 1 -ρ 1 + ρ 2 tan 2 φ 0 . (3)
The spectroscopic ellipsometer Ellipse 1891 developed in the Institute of Semiconductor physics of Siberian branch of Russian Academy of Sciences was used in the experiments. Measurements were performed at the angle of light incidence φ 0 = 70 over the photon energy range of 1.5-5.0 eV. The thickness of the modified surface layer d and information about the layer composition were obtained by fitting the measured data to model calculations. A multilayer model and the Bruggeman effective medium approximation (EMA) [START_REF] Bruggeman | [END_REF] are used. Selected model parameters are then varied (using the Levenburg-Marquardt algorithm) to minimize the error function. Two EMA models based on the XPS data have been used for cases of electrons and ions exposures, which will be considered in detail in corresponding Sections.

The reference data for the PD function of crystalline c-GaAs, amorphous a-GaAs and also for voids (empty spaces in a layer) were supplied together with the ellipsometer software. The GaAs native oxide and Ga 2 O 3 oxide data were taken from [9,10], respectively.

Electrodes of a gas-discharge system are known to be a subject of erosion under influence of ions and other highly-reactive discharge products, which lead to degradation of initially smooth electrode surface. Therefore, in the frame of the present investigation, it was of interest to track, whether the significant irradiation dose leads to deterioration of the surface of a semiconductor electrode.

A sample surface was inspected by means of optical microscope OPTEM with Mitutoyo objective (with lateral optical resolution of 1 µm) and atomic force microscope (AFM) INTEGRA Aura produced by NT-MDT company. The tapping mode with the average interaction force less than 1 nN and contact mode with average interaction force less than 100 nN were used for the surface profiling. Probes NSG01 and HA NC with the tip curvature radius of less than 35 nm, resonance frequencies of 150 and 120 kHz and force constants of 5 and 3.5 N/m, respectively, were used. The contact potential difference (CPD) distributions over a sample surface were measured by scanning two pass Kelvin probe microscopy (SKPM) mode. In the first pass, the surface topology was measured in the tapping mode. In the second pass the probe was lifted above the surface and repeated the recorded profile simultaneously measuring the CPD signal value between the AFM probe and the sample surface.

Experimental results

Exposure with electrons

In-depth profiles of chemical composition obtained with the XPS technique for an unprocessed GaAs sample are represented by Fig. 4. The data show relative concentration of gallium arsenide, gallium covalently bound to oxygen as well as arsenic bound to oxygen and for oxygen independent on the chemical bounds. All graphs are normalized to the sum of all measured components according to the equation (1), and the results are shown in atomic percent. One can see that the natural oxide layer (mixture Ga 2 O 3 and As 2 O 3 ) is present in the surface layer of the unprocessed sample. However, it is extremely thin: the thickness of the native oxide layer does not exceed one nanometer, while As 2 O 5 was not detected.

Figure 5 shows the concentration profiles for the surface layer of a sample exposed with electrons. The data refer to the exposure Q = 0.23 C/cm 2 . We also measured gallium and arsenic separately: The relative atomic concentrations of Ga and As in the sample bulk were close to 50%, which agrees with natural content of these elements in a GaAs crystal. Approaching the semiconductor surface, the concentration of GaAs decreases due to presence of the oxide film. Thickness of Ga 2 O 3 layer is about 5-7 nm. For the content of As 2 O 3 , a similar dependence is observed, whereas the As 2 O 5 content is negligibly small.

To investigate optical properties of GaAs as dependent on the electron irradiation dose, a set of samples exposed with doses in the range of 0.05-0.30 C/cm 2 was prepared. The result of their study with the spectroscopic ellipsometry technique is represented in Fig. 6, where the imaginary part of the PD function spectrum ε 2 is shown for different doses of exposure. The data refer to an untreated sample (a) and samples exposed by electrons with doses Q = 0.15 and 0.3 C/cm 2 , (b) and (c), correspondingly. The reference data for the imaginary part of PD function of crystalline GaAs is also shown.

The pseudo-dielectric function of a crystalline material is directly related to its electronic energy-band structure. So, the features observed in the ε 2 spectra are attributed to interband critical points of GaAs [11,12]. As seen in Fig. 6, the ellipsometric spectra contain peaks at E 1 , E 1 +∆ 1 (hν = 2.91 and 3.14 eV), E 0 (4.45 eV) and E 2 (4.77 eV) interband transitions, while contribution of other transitions such as E 0 and E 0 + ∆ 0 (hν=1.42, 1.75 eV) is weak and not observed clearly.

The most significant changes in the PD function spectrum for exposed samples take place in the 4.0-5.0 eV range. The amplitudes of E 0 and E 2 peaks become smaller compared to the case of the untreated surface, while the amplitude of the E 1 and E 1 +∆ 1 doublet do not change significantly. The higher the electron exposure dose, the lower are the E 0 and E 2 peaks amplitudes (data (b), (c)).

To obtain information about the chemical composition and structure of surface layers from SE data, a multilayer model was used. Usually, complicated structure of a surface layer, including its roughness and complexity in composition, is described with the effective medium approximation (EMA) model [START_REF] Bruggeman | [END_REF]. According to this approach, a real surface layer is replaced by a layer having a so-called effective medium dielectric function [START_REF] Azzam | Ellipsometry and Polarized Light[END_REF].

We point out that the ε 2 spectrum for the untreated sample (data (a)) is somewhat different from the corresponding reference spectrum of c-GaAs: namely, the amplitude of the peak at 4.75 eV is reduced essentially as compared to the c-GaAs curve. A natural suggestion could be that such a discrepancy is due to presence of both an oxide film on surface of the untreated sample, and of the transition layer between oxide and bulk semiconductor. (In this relation, see Fig. 4.)

An appropriate model for the analysis of ellipsometric data for a set of samples irradiated with electrons (including the untreated one) can be formulated on the base of XPS results. The oxide Ga 2 O 3 layer was found to appear on the GaAs substrate as a result of exposure with electrons (Fig. 5). This oxide is known [10] to have a larger refractive index compared to the native oxide [9], which consists typically of GaAsO 3 compound. So, to interpret experimental SE data for a set of samples irradiated with electrons, we developed a model that included GaAs substrate, the effective medium layer (which is a mixture of c-GaAs and the native oxide) and a thin film of native oxide.

The spectrum of ε 2 calculated for best-fitting parameters for an untreated sample is shown by solid line (a) in Fig. 6. Thus, the EMA layer of the thickness of 1.2 nm had a composition of 36.8% of GaAs and 63.2% of native oxide. The thickness of the native oxide film was evaluated to be 2.5 nm. (The oxide As 2 O 3 is not included into the model since its relative fraction is less than Ga 2 O 3 according to the XPS measurements.) Spectra of the imaginary part of the PD function calculated within such a model for samples exposed with different doses are shown by solid lines in Fig. 6. Apparently, there is a good correspondence of calculated and the experimental results. Parameters of the model applied in calculations are indicated in Table 1.

Results of fitting the calculated data to experimental curves, which are obtained for samples with different irradiation doses are summarized in Table 1. It shows that composition and thickness of the EMA layer depend on the irradiation dose: When it the thickness of the oxide layer Ga 2 O 3 grows. It reaches 4.5 nm maximal dose what is close to the data of Fig. 5. Figure 7 shows results obtained for the thickness of 2 O 3 and EMA layers as a function of dose.

Exposure with ions

As noticed in Section exposure of the samples with ions was carried out in the discharge system, when the semiconductor electrode operated at the negative polarity of the applied voltage. The XPS in-depth profiles of chemical composition for a sample subjected to the exposure by nitrogen ions with the maximal dose Q=0.47 C/cm 2 are represented by Fig. 8. The data show that there is no significant influence of irradiation with ions on chemical composition of a near-surface layer as compared to the untreated sample (Fig. 4). On the other hand, ellipsometric data demonstrate essential difference in optical properties of the untreated and exposed surfaces of the semiconductor. Figure 9 shows the imaginary part of PD function spectra ε 2 for samples exposed with ions at different doses. Data (a) are obtained for untreated surface and are the same as in Fig. 6. Data (b) and (c) correspond to spectra for surfaces exposed with ions of doses 0.23 and 0.47 C/cm 2 . One can see that there is a noticeable change in the ε 2 spectrum at an increase in the irradiation dose: The significant decrease in the spectra maxima and a slight red shift of the E 1 , E 1 + ∆ 1 doublet are observed, while the E 2 pike amplitude decreases also (cf. Fig. 6). When comparing these results with the corresponding data for the electron-exposed samples, one can see the essential difference in the behaviour of the PD function spectra for the two cases.

It is worth noting that similar changes in spectra of the PD function of GaAs samples are observed when the semiconductor surface is processed by ions during electron cyclotron resonance plasma etching [13], or at a low-energy hydrogen ion bombardment [14]. The ion implantation of elements into GaAs is known to be also accompanied by changes in the optical properties of semiconductor [15]. These effects are related to a thin film of damaged crystal structure, which is formed on the surface under ion bombardment. When an energetic ion penetrates a solid target, lattice atoms displacement occurs due to collisions between the incident ions and the near-surface atoms. In semiconductors, such displacement processes cause accumulation of structural defects within the lattice. Hence, an amorphous surface layer may be built up. As mentioned above, XPS measurements have not revealed essential changes in the chemical composition of the semiconductor due to its irradiation with ions. Therefore, we assume from the aforesaid that ion bombardment causes in our experiments only structural changes of the semiconductor surface, leaving the chemical composition of semiconductor material almost unchanged.

A usual way of describing optical response of such a material is to assume that the surface layer of the semiconductor is a physical mixture of crystalline and amorphous regions, and to use the EMA method to calculate the resulting PD function of modified GaAs [16]. However, in spite of its physical simplicity, the EMA model is not always able to describe the behaviour of PD function spectrum properly in the overall range of transition from the crystalline material to amorphous. As the authors of [15] have demonstrated, a better agreement between experimental and theoretical results can be achieved in the harmonic oscillator approximation (HOA), where the PD function is represented as a superposition of harmonic oscillators.

Nevertheless, as a first approximation to analyse the effect of irradiation of GaAs 2.

surface with ions, we restricted our consideration to the EMA model. The physical model consisted from the bulk semiconductor substrate combined with a modified GaAs layer. We did not include the oxide layer in the model, because the XPS measurements had not revealed the appearance of the oxide due to exposure with ions. The modified GaAs layer is assumed to be the mixture of crystalline c-GaAs, amorphous a-GaAs and voids. The spectra of imaginary part of PD function calculated with this model are shown by solid lines in Fig. 9 (b), (c), see the corresponding best-fitting parameters for these curves in Table 2.

One can see that the applied EMA model describes main features of spectra modification at the exposure with ions: A decrease in amplitude for both the doublet structure at photon energies near 3 eV, and of the pike at ∼5 eV. However, there are observed discrepancies between the results of the model and experimental data. At first, an expressed difference exists between experimental and theoretical curves in the range of low photon energy 1.5-2.5 eV. Furthermore, the model does not take into account the experimentally observed slight red shift of E 1 , E 1 + ∆ 1 doublet. As it has been pointed out above, more sophisticated models -like the harmonic oscillators approximationshould be used for a better consistency in theoretical description of the PD function spectra.

Parameters of modified surface model obtained for samples exposed by ions with doses in the range of Q=0.11-0.47 C/cm 2 are summarized in Table 2. According to these data, the ion irradiation leads formation structurally layer on the semiconductor It is interesting note that the experimental data can be described in the approximation, where the composition of this layer does not depend on the irradiation dose and is approximately: c-GaAs 37%; a-GaAs 44%; voids 19%. However, the thickness of the layer increases with Q, and the maximal measured thickness is ∼18 nm. The variation of the layer thickness, which is modified with the ions exposure, is shown in Fig. 10 for different exposure doses. The data for the untreated surface is shown in the figure also. We point out that colour of the sample surface changes after processing by large doses of particle flows at both discharge polarities. However, no erosion of the semiconductor surface was detected by means of optical microscopy. Further investigation of the surface quality of the samples was done with the atomic force microscopy.

AFM analysis of the semiconductor electrode surface.

An example of the AFM image of the surface exposed by electrons with dose of 0.37 C/cm 2 is shown in Fig. 11. The scanning area is 1 µm×1 µm. The average height of surface irregularities is about 1.5 nm; the distributions of the heights relatively to the average value is plotted at the bottom of the figure for the untreated sample (curve 1), and samples exposed with ions (2) and with electrons (3). These data clearly demonstrate that there is no significant difference between these distributions (the width variation is less than 6 Å). Increasing the scanned area to 10 µm×10 µm is accompanied by an increase of the characteristic width of height distribution; however, this effect is relatively small. We can conclude, therefore, that the discharge plasma treatment with the doses used in the experiments does not cause any noticeable erosion of the GaAs surface.

The contact potential difference (CPD) distribution over the surface area of the sample exposed by electrons with the dose of 0.37 C/cm 2 is shown in Fig. 12. The measurements were peformed in the two-pass mode (see Section 2). The distance between the Kelvin probe and semiconductor surface was 50 nm. The scanned area of 100 µm×100 µm was chosen in such a way that both untreated and exposed parts of the sample surface were investigated. The left side of the figure corresponds to the untreated surface. It is seen that there is a difference between CPD signals in both areas, which is about 300 mV. The CPD signal profile averaged over the area outlined by the white rectangle in the top part of the figure is plotted in the graph below. We note that the observed effect corresponds to the bending down the surface electron band of GaAs (at approximately 300 mV) compared to the initial state of the semiconductor. Similar results were obtained for the GaAs surface exposed with ions. At a comparable dose of the exposure, the absolute value of the CPD signal was approximately two times lower than for the exposition with electrons, while the sign of the effect is opposite.

Conclusion

Our results demonstrate that properties of a GaAs electrode are changed in the Townsend discharge due to exposure to the charge-carrier flux. The effect is observed while the semiconductor is exposed for both electrons and ions even for low current densities. The exposure of GaAs to electron flux stimulates growth of an oxide film on the semiconductor surface. There may be two possible sources of oxygen needed for the oxide growth in the gas-discharge environment: (i) Oxygen could be released from the SnO 2 film (which served as the second electrode to the discharge plasma) by its degradation due to exposure to active discharge products; (ii) It may be present as an impurity of a tiny content in nitrogen. We note that the oxide film also grows if pure (99.997%) nitrogen was used instead of nitrogen evaporated from liquid state.

Concerning the observed electron-induced acceleration of oxide growth on GaAs surface we note that in [17] oxide growth on silicon surfaces in presence of free electrons was reported. Although experiments in [17] were done for a different sample material and at different experimental conditions, we note that same processes (e.g. electron attachment to oxygen [START_REF] Yu | Gas Discharge Physics[END_REF]) may also activate the anode oxidation in our experiments.

When semiconductor serves as the cathode to the low-current Townsend discharge, no changes in the chemical composition of the semiconductor surface was detected with the XPS method. However, these data, together with results of spectroscopic ellipsometry measurements and measurements of surface potential with the AFM scanning Kelvin probe technique reveal essential structural changes in the GaAs surface induced by ion exposure. The results obtained give an evidence of partial amorphization of the GaAs surface layer, which thickness increases with the dose of ions.

We note that the average kinetic energy of ions impacting the cathode is rather low in our experiments and does not exceed several tenths of eV, see [START_REF] Brown | Basic Data of Plasma Physics[END_REF] for reference. The neutralization of such low-energy ions on a surface of metals and semiconductors is known to occur via the Auger neutralization process [START_REF] Hagstrum | [END_REF]: This occurs via tunnelling of electron from a deep electronic state of a solid to the ground state of ion that approaches its surface. The energy released at such a process, is transferred to the electronic subsystem of semiconductor. The amount of the transferred energy is rather large: It is in the order of ionization potential of the gas atom/molecule. In the course of relaxation of such excitations in the crystal, structural defects in the semiconductor lattice may be created near the semiconductor surface and accumulated there at subsequent ion neutralization events. It is essential that this mechanism of excitation of a semiconductor remains effective even at quite a low kinetic energy of ions bombarding the semiconductor surface.

In our experiments, plasma treatment of GaAs changes the surface potential of semiconductor, which should influence the electron work function and, correspondingly, the secondary electron emission coefficient γ [START_REF] Yu | Gas Discharge Physics[END_REF]. This effect may obviously be a reason for the drift of parameters of microdischarge devices (in which a semiconductor electrode is used) in time.

We also note, finally, that controllable modification of a semiconductor properties by a low-current Townsend discharge may be a practically interesting technological application of microdischarges.

Figure 1 .

 1 Figure 1. Experimental setup. 1 -GaAs sample; 2 -semitransparent to visible light nickel electrical contact; 3 -SnO 2 layer; 4 -glass plate; 5 -copper contact ring; 6dielectric spacer. Φ denotes visible and IR radiation illuminating the GaAs sample for the conductivity control. R is a 100 Ω resistor, V 0 -DC voltage source (the polarity was changed if the GaAs sample was exposed with ions).

Figure 2 .

 2 Figure 2. XPS-spectra of Ga 3d measured for irradiation dose Q = 0.23 C/cm 2 after different sputter times. Numbers at the right side is the sputter time in seconds. The transition from oxide (binding energy 20.2 eV) to GaAs (binding energy 19.0 eV) is obvious. The semiconductor served as the positive electrode of the gas-discharge system.

Figure 3 .

 3 Figure 3. XPS-spectra of As 3d measured for irradiation dose Q = 0.23 C/cm 2 after different sputter times. Numbers at the right side is the sputter time in seconds. The transition from arsenic oxide (binding energy 44.5 eV) to GaAs (binding energy 41 eV) can be proved by the chemical shift of the As 3d peak. The semiconductor served as the positive electrode of the gas-discharge system.

Figure 4 .

 4 Figure 4. XPS in-depth profiles of composition for unprocessed GaAs sample.

Figure 5 .

 5 Figure 5. XPS in-depth profiles of composition for GaAs sample irradiated by electrons with dose of 0.23 C/cm 2 .

Figure 6 .

 6 Figure 6. Imaginary part of pseudo-dielectric function spectrum obtained from SE data for samples irradiated with electrons with different doses Q. (a) untreated sample; (b) Q = 0.15 C/cm 2 ; (c) Q = 0.3 C/cm 2 .

Figure 7 .

 7 Figure 7. Thickness of Ga 2 O 3 oxide and EMA layer as a function of irradiation dose with electrons.

Figure 9 .

 9 Figure 9. Imaginary part of the PD function spectrum obtained from SE data for GaAs samples irradiated by ions with different doses Q. (a) untreated sample; (b) Q = 0.23 C/cm 2 ; (c) Q = 0.47 C/cm 2 . Points are experimental data, continuous curves calculation with the EMA model.

Figure 10 .

 10 Figure 10. Thickness of the modified GaAs layer as a function of irradiation dose with ions. The EMA model of the modified GaAs layer consists of a mixture of crystalline c-GaAs, amorphous a-GaAs and voids, see Table 2.

Figure 11 .

 11 Figure 11. Upper panel: AFM image of the surface exposed by electrons with dose of 0.37 C/cm 2 , see the text. Shown at the lower panel are distributions of amplitudes of surface irregularities for: (1) untreated surface; (2) surface exposed by ions with dose of 0.37 C/cm 2 ; (3) surface exposed by electrons with dose of 0.37 C/cm 2 . ∆h is the height deviation relative to the average value.

Figure 12 .

 12 Figure 12. Map of surface potential (upper part) and the potential profile for a sample exposed by electrons at dose of 0.37 C/cm 2 . The exposed part of the sample (right) is characterized by a lower potential.

Table 1 .

 1 Best-fitting parameters of the model used in processing ellipsometric data for samples irradiated by electrons.

	dose,		EMA layer		Ga 2 O 3 layer
	C/cm 2 GaAs, % Ga 2 O 3 , % thickness, nm thickness, nm
	0.05	32.7	67.3	2.7	1.0
	0.11	31.5	68.5	2.8	1.8
	0.15	25.3	74.7	3.2	2.3
	0.21	29.9	70.1	3.9	2.6
	0.26	24.1	75.9	3.3	3.6
	0.30	23	77	3.6	4.5

Table 2 .

 2 Best-fitting parameters of the model used in processing ellipsometric data for samples irradiated by ions

	dose,		composition of the modified layer
	C/cm 2 c-GaAs, % a-GaAs, % voids, % thickness, nm
	0.11	34	46	20	11.7
	0.23	36.8	44.2	19	16.3
	0.47	36.8	44.2	19	17.7

Acknowledgement

The work is partly supported by the President Grant of Russia for Scientific Schools (Grant 3306.2010.2), Russian Foundation for Basic Research (Grant 07-2-00168-a) and Program of fundamental research of Russian Academy of Sciences (Program No. 27).


Kaiyun Jiang
email: jiang@mch.rwth-aachen.de

Kostas Sarakinos

Stephanos Konstantinidis

Jochen M Schneider

Low temperature synthesis of α-Al 2 O 3 films by high power plasma assisted chemical vapor deposition

Keywords: 68.55.Nq, 68.60.Bs, 34.80.Dp -Al 2 O 3, PACVD, phase formation

Low temperature synthesis of -Al2O3 films by high-power plasma-assisted chemical vapour deposition

Introduction

Among the various Al 2 O 3 (alumina) polymorphs, the thermodynamically stable -Al 2 O 3 phase exhibits exceptional physical properties such as high hardness [1], transparency over a wide range of wavelengths [START_REF] Aguilar-Frutis | [END_REF]3], chemical inertness and thermal stability [1,3,4]. Due to these properties, -Al 2 O 3 films are widely employed in surface protection and microelectronics applications. In the industrial practice, -Al 2 O 3 is grown by thermal chemical vapor deposition (CVD). A common feature of CVD grown -Al 2 O 3 is that substrate temperatures in the excess of 1000 °C are necessary to provide the film forming species with the energy required for -Al 2 O 3 phase formation [5]. These high deposition temperatures may cause the build up of thermal stresses as well as interface reactions [5] and may thus lead to poor adhesion, and thermally induced crack formation and growth [6]. Furthermore, the metallurgical properties of the substrate material may be altered by substrate heating, severely limiting the choice of substrate materials. An alternative way to provide energy to the growing film and decreasing the temperature limit for the deposition of -Al 2 O 3 is by bombardment of energetic species which are available in plasma based physical and chemical vapor deposition techniques. For instance, Wallin et al. [7] demonstrated the deposition of -Al 2 O 3 on hard metal substrates at a temperature as low as 650 °C using high power pulsed magnetron sputtering, while Takamura et al. [START_REF] Yamada-Takamura | µm 1 µm (a)[END_REF] obtained this phase at a temperature of ~500 °C employing filtered cathodic arc. Jin et al. [9] reduced the temperature limit further for the growth of -Al In all the above cases, and despite the substantial decrease of the temperature limit for the -Al 2 O 3 formation, the deposited films have been found to exhibit porosity [11,12] and hence inferior elastic properties [12] compared to bulk and CVD -Al 2 O 3 [13]. In the case of films deposited by bipolar pulsed PACVD, Snyders et al. [11] suggested that the porosity is caused by chlorine (Cl) incorporation into the films, due to the insufficient disassociation of the AlCl 3 precursor employed during PACVD. It is known that an increase of the plasma density can also intensify the interaction of the plasma with the gas molecules and the growing film. The plasma density in turn depends on the power applied to the discharge. The PACVD discharges available so far have been operating at voltages of up to 0.9 kV [11,12,14], resulting in discharge power densities in the range 2 to 5 W/cm 2 . In the current study we have utilized a pulsed plasma generator which is able to deliver voltages up to 1.4 kV allowing for approx. 4 times larger power densities than those achieved in the conventional PACVD processes. Using plasma characterization tools and plasma modeling we show that increasing the discharge power density values from 5 W/cm 2 to 19 W/cm 2 results in an increase of the energy and the flux of the bombarding species towards the growing film, as well as in a more efficient AlCl 3 dissociation. These conditions allow for the growth of dense and phase pure -Al 2 O 3 films with negligible Cl incorporation and elastic properties close to the bulk values at a temperature of 560 ± 10 °C.

Experimental

Films were deposited in a semi-industrial parallel plate PACVD chamber from PLATEG GmbH.

The plasma was generated by applying unipolar voltage pulses (up to 1.4 kV in amplitude) with a nominal width in the range of 80 to 100 µs and a frequency of 5 kHz on two 16 cm in diameter steel electrodes. The substrates (Si (100) wafers) were placed on the cathode, while the other electrode (anode) was grounded. The power was supplied by a pulse generator developed by Ganciu et al. at Materia Nova R&D centre [15,16]. The experiments were performed in an AlCl 3 -Ar-H 2 -O 2 atmosphere at a pressure of 175 Pa. The composition of the gas atmosphere was 1.2% AlCl 3 , 14.8% Ar, 82.6% H 2 and 1.4% O 2 . The substrate temperature was kept at 560 °C using an external wall heating unit. The time dependent voltage and current signals were measured using an ELDITEST GE 8115 voltage and Tektronik A6303 current probe, respectively and recorded with a Tektronik TDS 3014B digital oscilloscope. To investigate the effect of the deposition conditions on the plasma composition, in-situ optical emission spectroscopy (OES) was employed. The OES measurements were performed using a PLASUS apparatus equipped with a charged couple device (CCD) camera for fast spectrum acquisition. The optical fiber was assembled together with a collimating lens to an observation window located at the chamber wall 10 mm above the cathode. The emission lines intensities corresponding to Ar (750.4 nm), H (434.0 nm), Al (396.1 and 394.4 nm), and AlCl (261.4 nm) neutral species and Ar + ions (476.4 nm) were sampled through a quartz window. The phase composition of the Al 2 O 3 films was determined by means of grazing incidence X-ray diffractometry (GIXRD) in a SIEMENS D5000 apparatus using a step size of 0.01°and an angle of incidence of 1°. Energy dispersive X-ray spectroscopy (EDX) measurements were carried out to determine the chemical composition of the deposited films. Scanning electron microscopy (SEM) was employed to study the influence of the deposition parameters on the film morphology, topography and thickness. Both EDX and SEM measurements were performed in a JEOL JSM-6480 apparatus, while for the EDX measurements an EDAX Genesis2000 detector was used. The elastic modulus and hardness of the as deposited films were measured by nanoindentation. The measurements were carried out using a Berkovich tip in a Hysitron nanoindentation system. The maximum load was 2500 µN to keep the contact depth 80 nm. The film thickness was ~1.0 µm.

The recorded data was analyzed using the method of Oliver and Pharr [17]. The increase of the voltage results in a close to linear increase of the current in the voltage range from 0.5 to 1.4 kV. The vertical dotted line (insert in figure 2) separates previously studied currentvoltage-range [10,11,18] from the range investigated here. The extension of the voltage range from 0.9 kV up to 1.4 kV results in power densities (calculated by dividing the power to the area of the cathode) up to 19 W/cm 2 , which are approx. 4 times larger than those reported previously [10,11,18]. The increase of the power density affects the plasma composition, as determined from the OES data shown in figure 2. The emission intensity of all lines increases as the discharge voltage and hence the cathode power density is increased. This may be consistent with an increase of the electron density and/or temperature, since the line intensity is proportional to the number density of excited atoms which in turn is related to the number density of the considered atoms in their ground states and to the density of the energetic particle responsible for the excitation. For instance, the energetic particle can be either an electron or an argon metastable atom (Ar*). As the voltage is increased, it is reasonable to assume that electrons will acquire more energy from the applied electric field resulting in more efficient dissociation and ionization of the plasma species. As reported in figure 2 Therefore, the increase of the emission intensity of precursor related Al and AlCl lines in figure 2 should not only be attributed to the electron temperature and density increase, but also to the enhancement of the dissociation by impact of Ar*. Enhanced dissociation of BCl 3 has also been reported upon increasing the discharge power in a PACVD plasma [20].

Results and discussion

Measurements

Apart from its effect on the plasma composition, the increase of the cathode voltage to 1.3 kV leads to an up to two times larger discharge current as compared to the 0.9 kV discharge, as shown in the insert in figure 2, increasing the ion flux impinging on the growing film surface. In addition, the increase of the voltage implies a larger potential drop across the cathode sheath which in turn has implications for the energy of the ions impinging on the film. Due to the relatively high pressure of the discharge (175 Pa), ion-neutral and ion-electron collisions occur in the sheath [21]. Hence, the average ion energy is lower than the energy that can be gained by acceleration across the cathode fall. The ion energy distribution function (IEDF) was estimated based on the Davis and Vanderslice model [21] considering charge exchange collisions in the sheath. The ion energy distribution function is then given by the following expression, [ ]

) U/U - 1 - (1 / - exp U/U - 1 2 1 dU N dN U / 2 c c o c ⋅ ⋅ = ⋅ ⋅ ⋅ λ λ L L (3)
where L/ is ratio between sheath thickness and mean free path for charge exchange collisions, N 0 is the total number of ions at the sheath edge and U c is the potential difference across the sheath, i.e. the cathode potential. Here we evaluated for simplicity an Ar-H 2 plasma with a pressure of 175 Pa, since Ar and H 2 constitutes ~97% of the gas atmosphere IEDFs were calculated using Eq. (3) for pure Ar and H 2 plasmas and plotted in Fig. 3 (solid lines). For the calculation of in Eq. (3), the cross sections for Ar-Ar + (3× 10 -15 cm 2 ) and H 2 -H 2 + (8× 10 -16 cm 2 ) charge-exchange collisions were employed for Ar and H 2 plasma, respectively [22,23]. The sheath thickness L was estimated based on the Child-Langmuir law to be 1 mm and 8 mm for Ar plasma and H 2 plasma, respectively [24].

The dashed curve in figure 3 corresponds to a plasma that consists 85% of H 2 and 15% of Ar and obtained by averaging the data from the H 2 and Ar plasma IEDFs using the respective weighting coefficients 0.85 and 0.15. These conditions are close to our experimental plasma conditions and show that the IEDF exhibits a low populated high energy tail starting from energies that correspond to a value 1/3U c , while the majority of the species exhibit energies that correspond to values below 1/4U c . Furthermore, the average ion energy determined by integration of the IEDF gives a value=0.16 U c . Therefore, as U c is increased from 0.9 to 1.3 kV the mean average ion energy increases from ~140 to ~200 eV. The increase of the energy and the flux of the ions towards the growing film surface result in a more intense energetic bombardment. formation and improvement of crystallinity thereof can be attributed to the increased average ion energy as discussed in the pervious paragraph as well as to the increase in duration of ion bombardment enabling surface diffusion [10,13,18,25]. It has to be mentioned here that due to the high plasma density at the voltage and pulse width of 1.3 kV and 100 µs, respectively, the substrate temperature was 590 ± 10 °C although no external heating was applied. The deposition temperature for all other films was 560 ± 10 °C.

Θ ~37.8°2 Θ ~43.3°2 Θ ~52.5°2 Θ ~25.6°2 Θ ~37.8°2 Θ ~43.3°2 Θ ~52.5°F WHM (°) voltage (kV)
Figure 5. FWHM of chosen diffraction peaks as a function of various discharge voltages and pulse widths. Due to difficulty of peak fitting, data from sample deposited at 0.9 kV/80µs at 2 ~25.6°is not shown.

Along with the change of the phase composition and the improvement of the crystallinity, a continuous decrease of the Cl incorporation from 1.1 to 0.5 at.% is observed as the voltage is increased from 0.9 to 1.3 kV. This may be understood by considering recently published Cl desorption data by Snyders et al. [11] where Cl release from -alumina films was measured during annealing at temperatures larger than the deposition temperature. This indicates that desorption of Cl may also be expected during deposition as the energy at the surface of the growing film is increased for example due to enhanced ion bombardment. Cl desorption during growth in turn results in a decreased Cl concentration in the film, which is again consistent with the experimental observation. Recently published ab-initio data [11] clearly indicated that the presence of Cl destabilizes alumina, which underlines the novelty of the here reported process where the Cl concentration could be decreased by an increase of the energy and the flux of the bombarding species towards the growing film, as well as in a more efficient AlCl 3 dissociation. As the discharge voltage is increased, the Cl incorporation is reduced and the formation of the -Al 2 O 3 phase is observed. Similar behavior is also obtained for the films grown at a pulse width of 100 µs. At the same time the deposition rate was found to decrease from 0.5 µm/h at a voltage of 0.9 kV to 0.36 µm/h at a voltage of 1.3 kV (pulse width of 80 µs in both cases) which could be associated with the increase of resputtering and density of deposited film at elevated voltage. On the other hand, the change of the pulse width from 80 to 100 µs had no measurable implications for the deposition rate.

The elastic modulus (E) and hardness (H) of the deposited films are plotted versus the voltage in figure 6. The quantities E and H increase as the voltage is increased. The highest values for E (~380 GPa) are obtained for the films synthesized at a voltage of 1.3 kV and are very similar to that of CVD grown -Al 2 O 3 [26] and bulk ceramic -Al 2 O 3 [27] as indicated by the horizontal dashed lines in figure 6. A top view (a) and a cross sectional (b) SEM micrographs of this film are presented in the figure 7a andb, respectively. These images show a smooth surface and a dense morphology, which can be understood by the formation of phase pure -Al 2 O 3 with a density close to the bulk value, which is a consequence of ion bombardment at elevated cathode voltage, which is characterized by both a larger flux and larger average ion energy as compared to conventional PACVD synthesis approaches.

Conclusions

In conclusion, we have employed a novel generator that allows for the operation of a PACVD discharge at approx. 4 times larger power density than those utilized conventionally. Plasma analysis and modeling indicate that the larger discharge power density enables a significant increase of AlCl 3 precursor dissociation efficiency, as well as an increased ion flux towards the growing film.

These conditions enable the deposition of smooth and dense -Al 2 O 3 films with negligible Cl incorporation and elastic properties similar to those of the bulk -Al 2 O 3 at a growth temperature of 560 ± 10 °C.

2 O 3 down

 23 to 280 °C by depositing Al 2 O 3 films on Cr 2 O 3 seed layer by rf magnetron sputtering. Kyrylov et al. [10] demonstrated the growth of -Al 2 O 3 at 560 °C by plasma assisted chemical vapor deposition (PACVD).

  of the time dependent current (I) and voltage (U) of a discharge operating at a pulse width of 80 µs and a frequency of 5 kHz are shown in figure 1. Both current and voltage exhibit constant values during the nominal pulse on-time (80 µs). As the pulse is switched off, current and voltage decay over approx. 70 µs towards zero. In the following text the magnitudes of current and voltage refer to the values measured during the pulse on-time. The I-U curve of a discharge operating at a pulse width of 80 µs and a frequency of 5 kHz is presented in the insert in figure 2.

Figure 1 .Figure 2 .

 12 Figure 1. Time dependent cathode voltage (bottom panel) and current (top panel) of a discharge operating at a pulse width of 80 µs and a frequency of 5 kHz.

Figure 3 .

 3 Figure 3. Ion energy distribution functions for pure Ar and H 2 plasmas (solid lines) and a plasma consisting of 85% H 2 and 15% Ar (dashed line).In figure4, the XRD patterns of films grown at various cathode potentials are presented. At 0.9 kV cathode potential the formation of a mixture of γ-and -Al 2 O 3 is observed. As the voltage is increased to 1.1 kV, phase pure -Al 2 O 3 films are obtained. The increase of the voltage enhances the crystallinity of the -Al 2 O 3 phase as manifested by the decrease of the full width at half maximum (FWHM) of the XRD peaks. As shown in figure5, FWHMs of peaks at 2 angular positions of 25.6°, 37.8°43.3°and 52.5°, obtained by Gaussian fitting, show a decrease by 30-50%, when discharge voltage increases from 0.9 to 1.3 kV. Meanwhile, a clear appearance/separation of peaks at 2 positions of 35.1°and 68.2°is also observed in figure4at a voltage of 1.1 and 1.3 kV, also indicating improved crystal quality. Furthermore, an increase of the pulse width from 80 to 100 µs (demonstrated for voltages of 1.1 and 1.3 kV in figure4) leads to a decrease of the FHWM of the -Al 2 O 3 peaks by 5~10% (see figure5) and subsequently to an improved crystallinity. The -Al 2 O 3

Figure 4 .

 4 Figure 4. X-ray diffraction patterns of Al 2 O 3 films deposited at various discharge voltages and pulse widths. The inverse triangles and the circles indicate the positions of the diffraction peaks associated with the presence of -and γ-Al 2 O 3 , respectively.

Figure 6 .

 6 Figure 6. Elastic modulus (upper) and hardness (lower) of Al 2 O 3 films as a function of the discharge voltage. The insert is a top view and the cross section SEM micrograph from a film deposited at a discharge voltage of 1.3 kV and a pulse width of 100 µs.

Figure 1 .Figure 2 .

 12 Figure 1. Time dependent cathode voltage (bottom panel) and current (top panel) of a discharge operating at a pulse width of 80 µs and a frequency of 5 kHz.

Figure 3 .

 3 Figure 3. Ion energy distribution functions for pure Ar and H 2 plasmas (solid lines) and a plasma consisting of 85% H 2 and 15% Ar (dashed line).

Figure 4 .Figure 5 . 3 Figure 6 .Figure 7 .

 45367 Figure 4. X-ray diffraction patterns of Al 2 O 3 films deposited at various discharge voltages and pulse widths. The inverse triangles and the circles indicate the positions of the diffraction peaks associated with the presence of -and γ-Al 2 O 3 , respectively.

  , the line intensity of the Ar 750 nm line versus the applied voltage increases. The low energy level of this transition is related to the 3P 0 and 3P 2 metastable levels of the argon atom. Hence it is reasonable to assume that the increase of the line intensity indicates an increased population of metastable Ar (Ar*) as the discharge voltage is increased. Consequently, more Ar* would be available for dissociating the precursor. Zheng et al. reported that in an inductively coupled discharge, the presence of Ar* leads to efficient AlCl 3 precursor dissociation[19]. Generally, these dissociation processes can be described by the following reaction[19],

	AlCl	3	+	Ar	*	or	e	-	→	AlCl	+	2Cl	+	Ar	or	e	-	(1)
	which leads to the formation of Al through the reaction HCl Al H AlCl + → +							(2)

Acknowledgments

The authors acknowledge the funding from the Deutsche Forschungsgemeinschaft (DFG) within Collaborative Research Centre (TFB) 289. S. Konstantinidis is postdoctoral researcher at the National Fund for Scientific Research (FNRS, Belgium), and he thanks The Interuniversity Attraction Poles Program of the Belgian Science Policy (Project "PSI: Fundamentals of Plasma Surface Interactions") for the financial support.


P M P Salomé
email: psalome@ua.pt

J Malaquias
email: ajmalaquias@ua.pt

P A Fernandes
email: bpafernandes@ua.pt

F Da Cunha

A F Da Cunha

Mo bilayer for thin film photovoltaics revisited

Keywords: Mo, molybdenum, solar cells, thin films, electrical contact, bi-layer PACS: 68.37.-d, 68.55.-a, 81.15.Ef, 81.15.Cd, 73.61 -r

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Introduction

Thin film photovoltaics (PV) is an emerging worldwide industry. Due to the recent increase in energy prices and with more ecologic global concerns, PV has increasingly become a more important technology.

From all the current PV technologies, the one that is showing more industrial interest is the Cu(In,Ga)Se2 (CIGS) based photovoltaics. Several authors [1,2] believe that in the next decade, solar modules made from CIGS will overtake the giga Watt per year production. This type of technology is usually known as thin film PV, because, when compared with Si, the absorber layer is much thinner.

CIGS thin film solar cells are based on a layered structure represented in figure 1. Each one of the layers has a different role in working cells. The work presented in this report refers to the back contact layer. This layer has one main purpose which is basically to make the positive electrical contact of the cell. Although that is the main purpose, there are a few other requirements for this layer that need to be fulfilled, these are: have low resistivity; have good adhesion to soda lime glass (SLG); have low roughness; be chemically inert with the materials deposited on top, mainly Cu, In and Ga and Se; be stable during the high growth temperatures of the CIGS; allow Na diffusion from SLG since the properties of the CIGS films are influenced by the Na doping that occurs during growth, this has a positive effect in the solar cells performance [START_REF] Rudmann | Effects of sodium on growth and properties of Cu(In,Ga)Se2 thin films and solar cells[END_REF]; have a high reflectance in the visible spectrum; have a temperature expansion coefficient similar to SLG and CIGS and form an ohmic contact with CIGS. One material that suits almost all those requirements is molybdenum (Mo). Mo is used by a large number of research groups and by the CIGS industry already in the market as it is shown in table 1. As comparison, it is also shown the CIGS growth method and the substrate. Despite different methods and substrates used in all these companies, Mo as back contact is common to all of them. Mo is a metal which can be deposited as thin films with a good resistivity, a good adhesion and fairly smooth. It does not react with Cu, Ga or In and although it reacts with Se forming a MoSe 2 [4] layer there are small positive consequences for the solar cell performances reported [5]. Mo is also able to allow diffusion of Na [START_REF] Rudmann | Effects of sodium on growth and properties of Cu(In,Ga)Se2 thin films and solar cells[END_REF] during CIGS growth, and it withstands its growth conditions. One weak point that is usually acknowledged in using Mo, is its low reflectance in the visible part of the spectrum. Some groups [7] are trying to lower the thickness of the CIGS, their motivation is to make PV modules cheaper by using lower quantities of material. With smaller thickness in the absorber layer, the amount of light absorbed will decrease and therefore a good reflectance of light in the back contact is required. The basic idea is to increase the number of absorbed photons by reflecting those that are not absorbed, in the first passage, by the absorber.

Other metals with low resistivity and reflectivity values, for example Pt, or Au, usually diffuse into the CIGS layer during the deposition [8]. Other reports have confirmed that cells made with Mo perform better than cells with a back layer with better reflectivity than Mo [8].

Mo is usually deposited by sputtering, mainly because these methods are used in industry and are easy to work with. Although Mo appears to be a good material to be used as back contact, its deposition is not without problems. When depositing Mo by these methods, the resistivity and the adhesion of the films are opposite effects making it difficult to have a film with both properties maximized. The reason why this happens is usually attributed to the film stress [9,10] state. Films that are deposited under compressive stress have low resistivity but also they have low adhesion.

On the other hand, films that are deposited under tensile stress have good adhesion but high resistivity. To overcome this problem Scofield et al [10] suggested depositing a bi-layer. If we can have a first layer with good adhesion followed by a layer with good resistivity, it might be possible to have a layer in which both properties would be adequate.

Despite the fact that the properties of Mo are well known and several groups already use, for a long time, a Mo-bilayer, the clear deposition conditions of such a layer have not been reported so far. Therefore the objective of this work is not to study all the referred properties but to present suitable deposition conditions of a Mo bi-layer. So, from the listed requirements, only the study of the properties that can be changed by the deposition method will be done. The properties that are required to make suitable Mo films are: thicknesses no bigger than 1000 nm; sheet resistances lower than 1 ; good adhesion and being able to support temperatures up to 580 ºC and maintain the same properties as before.

Influence of the sputtering pressure

Sputtering working pressure influences the properties of sputtered atoms mainly by two ways [START_REF] Jonh | Sputter deposition Processes in Handbook of Deposition Technologies for films and coatings[END_REF][START_REF] Mattox | Handbook of Physical Vapor Deposition (PVD) Processing[END_REF]. The first is by changing the deposition rate and the second is the mean free path.

With more pressure in the chamber the probability of a sputtered atom colliding with an Ar atom increases and thus changing the energy of the sputtered atoms. The energy with which the sputtered atoms arrive at the substrate is an important factor that determines the stress state of the films.

When sputtering at low pressures, Mo films have the tendency to become tightly packed, this tends to form films with a compressive stress and there is a decrease in resistivity. On the other hand, when sputtering at high pressures Mo films have a more porous columnar grain growth causing intergranular voids. This happens due to the reduced energy of sputtered atoms that arrive at the substrate. In terms of film properties, these voids make the resistivity higher but since there is not a reorganization of the arriving atoms, the adhesion is good. The tensile forces between these grains are attractive and inversely proportional in strength to the intergranular spacing [9].

Although these two explanations seem to be followed by different groups, there is an effect that is not fully explained yet. It is reported that Mo films have a maximum value in tensile stress and with increasing deposition pressure the tensile stress lowers and even gets to compressive stress.

This effect is not explained so far.

The stress state of the films is then dependent on the microstructure of the Mo films which is also related with the deposition conditions. In ideal situations we would require to have the electrical behaviour of the compact films deposited at low pressures with compressive stress and the good adhesion that films with columnar grains show when they are in a tensile stress deposited at high pressures. The stress state of the films studied in this work will be estimated using the rocking curve method.

The Rocking Curve method

The rocking curve method is a non destructive x-ray based method that allows us to estimate the stress state of a thin film. It is a complex technique that explores the strain done by a stress produced in the film. Let us consider that , is the elastic lattice strain in a direction defined by the Euler angles φ and with respect to the specimen frame of reference, see Error! Reference source not found.figure 2. The strain depends on the components of the stress tensor

ij [ ij = ji (i j); i3 = 3i =0] [9,13] as: ( ) ψ σ σ σ ε φ φ ψ φ 2 2 22 11 1 0 0 , , sin 2
1 hkl hkl hkl hkl hkl hkl S S d d d + + = - = Ψ (1) with φ σ φ σ φ σ σ φ 2 22 12 2 11 sin 2 sin cos + + = (2)
Where is the angle between the normal to the diffracting set of lattice planes [hkl] and the normal to the specimen surface, φ is the rotation angle within the plane of the layer (see Error! Reference source not found.Figure 2). hkl d ψ φ, is the lattice spacing of [hkl] planes in the direction defined by φ and , and hkl d 0 is the corresponding strain-free lattice spacing of [hkl] planes. For the case of Mo, the direction used is [321]. The peak is found around the 2 range 129- 137 º.

Experimental

Film Deposition

Deposition of Mo was done by dc magnetron sputtering. The Mo target purity was N4 and it was used Ar N5 as the sputering gas. Mo films were deposited on 3×3 cm 2 soda lime glass (SLG), which were cleaned with acetone, alcohol, hot deionised water and dried using a N 2 flux.

The base pressure of the chamber was 5×10 -6 mbar. Thicknesses were monitored with a quartz crystal monitor Intellemetrics IL 150. The working pressure in the chamber could be changed in the interval of 1x10 -3 and 1x10 -2 mbar. This was done by changing the Ar flow or by opening a butterfly valve. In this study we have always used the same flow of Ar to calibrate the butterfly position at a predefined pressure. Pressure then could be changed using different Ar flow values.

We used a sputtering power of density of 0.5 W/cm 2 and the substrate-to-target distance was 10 cm.

Film Characterization

Several methods have been used to characterize the films. Four Point Probe was used to measure the Sheet Resistance of the films; a Dektak 150 step profiler was used to measure film thicknesses; reflectivity was measured using a double beam Shimadzu spectrophotometer UV-3600; films' adhesion was tested by the Scotch tape test [10] and the rocking curve method was used to calculate the stress state of the films. finally XRD analysis were used to see the preferential orientation of the Mo films. X-ray diffraction (XRD) was done with a PHILIPS PW 3710 with a Cu-K line of 1.5406 Å. Scanning electron microscopy (SEM) was used to see the morphology of the sample and a Hitachi SU-70 with a acceleration voltage of 7 kV was used.

It was considered that the Scotch tape test or adhesion test had only two results, good or bad. If after the application of the tape the film was not damaged, then the adhesion was good, otherwise it was considered bad.

Resistivity

As stated before, the four point probe method was used to measure the sheet resistance ( s ) of the films. Knowing the sheet resistance and the thickness (h) of the films it is possible to estimate the resistivity ( ) using the following equation:

s h ρ ρ ⋅ = (3)
This equation is valid for square uniform samples.

Rocking Curves

For these measurements a x-ray diffractometter with a CuK radiation was used. For measuring the profiles of the Mo [321] peak, a scan between 0º and 63º in the direction was done. We used Bragg's Law to calculate . The linearization was done in the interval of 0.25 to 0.7 sin 2 . 321 2 S takes the value of 4.76x10 -6 MPa -1 [9,10] and therefore it is possible to calculate φ σ . Although, it is also possible to determine bi-axial components of the stress tensor, since these follow the general behaviour of φ σ , this component is enough to evaluate the stress state of the films.

Results and Discussions

Single layer films

Single layer Mo films were initially deposited to test the different regimes described by Scofield et al [START_REF] Jonh | Sputter deposition Processes in Handbook of Deposition Technologies for films and coatings[END_REF] and analyse the reproducibility of such conditions in our chamber. In table 2 the results of the first experiments are shown. For this set of experiments the main purpose was to identify the changes in adhesion and resistivity as function of the pressure.

Table 2 shows the same tendency that has been reported by Scofield et al [10] for all properties. Films sputtered at high pressure have good adhesion and bad resistivity and films at lower pressure have the opposite properties. In Figure 3 it is possible to see the resistivity as function of the pressure. The results follow the same tendency as shown in previous studies [9,10]. From these results it is evident that there are two regimes when depositing Mo. The first regime, at pressures between 6x10 -3 and 10x10 -3 mbar, produces samples with good adhesion but with a high resistivity. In the second regime, at pressures between 4x10 -3 and 1x10 -3 mbar, the samples have bad adhesion but a lower resistivity as it is shown in Figure 3.

This analysis demonstrates what regimes we can use to do the bi-layer. The first layer should be deposited in a pressure higher than 6x10 -3 mbar to obtain good adhesion and the second with a pressure smaller then 4x10 -3 mbar to obtain a low resistivity.

Bi-layer films

For this test it was decided to deposit a layer at high pressure at the bottom followed by a layer at low pressure at the top. It was decided to make the bottom layer with 500 nm thick and to study the variation of the thickness of the top layer. The pressures chosen were 10x10 -3 mbar for to the bottom layer and 4x10 -3 mbar for to the top layer. Between the two different pressures there is a period of 2 minutes. During this time the pressure is lowered in 1 minute and the tooling factor is changed in the other minute. Since this is not automatic these transitions are different from film to film and can change the final thickness of the film. By looking at the properties presented in Table 3, we see that when the top layer is thin, then the film properties are close to those of the single layer samples deposited at high pressure. In the other extreme case, where the top layer is as thick as the bottom layer, the film has bad adhesion and the lowest resistivity. The best sample of this run was the one in which the bottom layer had 500 nm thick and the top layer had 300 nm thick. For the top layer of 400 nm there was poor adhesion and so continuing to increase the thickness of the top layer was useless. The value of 1.8 for the sheet resistance is not low enough and therefore the deposition pressure of the top layer was further decreased.

In Table 4 the properties of the Mo films with a top layer deposited at 1x10 -3 mbar are presented. Four different second layers were used, 50 nm, 200 nm, 300 nm and 450 nm. This time with increasing thickness of the top layer no noticeable degradation in the adhesion was observed.

In the case in which the top layer had 450 nm thick it was noticed a degradation in the value of the resistivity as shown in Figure 4. This was due to a poorer adhesion but since the tape test only dictates a fail or pass evaluation, the degree of adhesion is not quantified. If the adhesion of film is degraded, then the morphology of the film is degraded and consequently resistivity is increased. Sample number 3 presented the lowest resistivity and had a good adhesion. This sample showed a sheet resistance of 0.6 with a total thickness of 817 nm which corresponds to a resistivity of 49.1 µ .cm. This value of resistivity is very close to that of a Mo single layer deposited at 1x10 -3 mbar. There are some differences between the monitored thicknesses and the ones measured with the profiler but these differences are small and acceptable. This value is in accordance with values reported for other works [14,15] Since it was shown for these pressure ranges that the lowest achievable resistivity was 49.1

µ cm for the sample with thickness 500+300 nm at 10x10 -3 /1x10 -3 mbar, it was decided to study further samples deposited in these conditions.

Characterization of the bi-layer films

An important question that has arisen from the last test was the reproducibility. Since there was a difference between the crystal monitor values and the measured thickness values it was important to check if this error was systematic or random. To investigate this problem, several samples were deposited. This was done for more than 4 runs, each with 4 films. It was then noticed that all the films had good adhesion. Their thicknesses were between 800-900 nm and the sheet resistance values were between 0.5 and 0.8 . These differences were most likely due to the pressure transition time between the bottom and top layers. This transition was done manually and therefore variations may exist from sample to sample.

To ensure that the film could resist the temperature of growth of CIGS, one of the samples was placed in the deposition chamber and heated up to 580 ºC for one hour. The Mo film was able to sustain these high temperatures.

The XRD spectrum presented in Figure 5 Calculations of the stress tensor, φ σ , were made as previously explained. As shown in Figure 6 b), not only the correlation between the points is low, but the slope of the curve is also low. Although the correlation is weak (R 2 of 0.7), it is evident that there is a general tendency for higher values of 321 ,Ψ φ d with the increase of sin 2 .This means that the film is in a tensile state. The result of the stress tensor component, φ σ , is 17 MPa, which is a low value. For comparison purposes, the values found by Vink et al [9] were in the order of GPa. This difference is possibly due to the fact that the film is too thick, but it is also possible that this is a mean value for a film with the top half in a compressive state and the bottom half in a tensile state.

Nevertheless the important film properties had the desirable values.

In figure 7, a SEM image of the cross section of a Mo bi-layer is presented. The film shows a columnar growth and it is possible to see the bottom layer with 500nm and the top layer with 300 nm.

Conclusions and Outlook

The main goal of this work was to clearly identify the deposition conditions, lacking in the literature, that ensured the deposition of Mo bi-layers with good adhesion to the SLG, a sheet resistivity lower than 1 and good reproducibility of the former two properties. In this paper we presented the results of a study for varying pressure and thickness of the top layer. A detailed study of the resulting bi-layers was made, and the optimal result was achieved when the bottom layer was deposited at 10x10 -3 mbar with a thickness of 500 nm and the top layer deposited at 1x10 -3 mbar with a thickness of 300 nm. The bi-layers produced by this method meet the desired properties by having a good adhesion and a sheet resistivity between 0.5 and 0.8 .

The resulting bi-layers were analysed and XRD studies showed a film with the typical Mo cubic structure and a crystallite size of 40 nm. Stress calculations showed that these films were in a low tensile state. These calculations were difficult to make due to the big dispersion of the data when analysing the rocking curves itself. The method proved to be reproducible and the films could withstand the harsh growth conditions of CIGS.

Figure 1 :

 1 Figure 1: Layer sequence of a CIGS solar cell device, SLG stands for soda lime glass and TCO for transparent conductive oxide. The figure is not at scale.

Figure 2 :

 2 Figure 2: Definition of Euler angles and with respect to the overall stress state.

  elastic constants for lattice plane[hkl]. For a polycrystalline specimen with random crystallite orientation and a negligible variation of ij over the depth probed, hkl ψ φ ε , depends linearly on sin 2 .

2 )

 2 and equation 1 it is possible to estimate the value

Figure 3 :

 3 Figure 3: Mo film resistivity versus working pressure.

Figure 4 :

 4 Figure 4: Resistivity versus thickness for samples deposited at with a bottom layer of 500 nm deposited at 10x10 -3 mbar and different top layers deposited at 1x10 -3 mbar.

Figure 5 :

 5 Figure 5: XRD spectrum of a Mo bi-layer film.

Figure 6 :

 6 Figure 6: a) Example of Voigt fitting for = 35,84; b) 321 ,Ψ φ d

Figure 7 :

 7 Figure 7: SEM image of the cross section of a Mo-bilayer.

Table 1 :

 1 Substrates, CIGS growth methods and back contact material for several CIGS companies[6].

	Company	Substrate	Back Contact	Growth CIGS Process
	Shell Solar	Glass	Mo	Sputter/Selenization
	Global Solar	Steel	Mo	Coevaporation
	Miasole	Glass	Mo	Sputter/Selenization
	Würth Solar	Glass	Mo	Coevaporation
	Avancis	Glass	Mo	Sputter/RTP
	Daystar Tech	Glass	Mo	Sputter
	EPV	Glass	Mo	Sputter/evaporation
	Ascent Solar	Polymer	Mo	Coevaporation
	ISET	Glass/Flex	Mo	Ink/Selenization
	Nanosolar	Flexible	Mo	Nanoink print/ RTP
	Heliovolt	Glass	Mo	FASST
	SoloPower	Steel	Mo	ED/RTP

Table 2 .

 2 Properties of single layer Mo films.

	Pressure (mbar)	Adhesion	Resistivity (µ .cm)
	10x10 -3	Good	411.0
	6x10 -3	Good	330.4
	4x10 -3	Bad	82.7
	1x10 -3	Bad	36.3

Table 3 :

 3 Description of the bi-layer process with 10x10 -3 /4x10 -3 mbar.

	Sample Step	Pressure (mbar)	Monitor Thickness (nm)	Total (nm) measured thickness	Adhesion	s ( ) (µ .cm)
		1	10x10 -3	500
	1	2	4x10 -3	200	685	Good	4.5	307.68
		1	10x10 -3	500
	2	2	4x10 -3	300	808	Good	2.4	194.07
	3	1 2	10x10 -3 4x10 -3	500 400	911	Bad	1.8	163.98

Table 4 :

 4 Parameters and results of the bi-layer deposition with 10x10 -3 /1x10 -3 mbar pressures.

	Sample Step	Pressure (mbar)	Controller Thickness (nm)	Adhesion	Total (nm) measurement thickness	s ( ) (µ .cm)
		1	10x10 -3	500
	1	2	1x10 -3	50	Good	567	1.6	90.7
		1	10x10 -3	500
	2	2	1x10 -3	200	Good	717	0.8	57.4
	3	1 2	10x10 -3 1x10 -3	500 300	Good	817	0.6	49.1
	4	1 2	10x10 -3 1x10 -3	500 450	Good	972	0.8	77.8

Acknowledgments P.M.P. Salomé acknowledges the financial support of Fundação para a Ciência e Tecnologia (FCT), through a PhD grant number SFRH/BD/29881/2006. P. A. Fernandes thanks the financial support of the FCT, through a PhD grant number SFRH/BD/49220/2008. The national network of microscopy is acknowledged: FCT: REDE/1509/RME/2005.

The authors also acknowledge Eng.ª Rosario Soares for the rocking curves measurements.