Light-induced changes in the structure and optical dispersion and absorption of amorphous As40S20Se40 thin films

Materials Chemistry and Physics 115 (2009) 751–756

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Light-induced changes in the structure and optical dispersion and absorption of amorphous As40 S20 Se40 thin films E. Márquez ∗ , R. Jiménez-Garay, J.M. González-Leal Department of Condensed Matter Physics, Faculty of Sciences, University of Cadiz, 11510 – Puerto Real, Andalusia, Spain

a r t i c l e

i n f o

Article history: Received 26 September 2008 Received in revised form 24 October 2008 Accepted 16 February 2009 Keywords: Chalcogenide glasses Thin films Optical properties Photo-induced effects

a b s t r a c t Exposure with bandgap light, in air, and thermal annealing at a temperature near the glass transition temperature, of thermally-evaporated amorphous As40 S20 Se40 thin films, were found to be accompanied by structural effects, which, in turn, lead to changes in the refractive index and shifts in the optical absorption edge. Also, clear indications of photo-oxidation were found after light exposure, in air. The dispersion of the refractive index is discussed in terms of the single-oscillator Wemple–DiDomenico model. The strongabsorption region of the absorption edge is described using the ‘non-direct electronic transition’ model, proposed by Tauc. Regarding the structural transformations that take place in As40 S20 Se40 chalcogenide thin films, when exposed or annealed, changes in the first sharp diffraction peak, present in the X-ray diffraction pattern, with both treatments, has been interpreted as a diminution of the interstitial volume around the AsS3−n Sen (n = 0,1,2,3) pyramidal structural units, which form the amorphous network of the samples under study. This result is certainly consistent with the decrease of the average thickness found for the illuminated and annealed chalcogenide samples. © 2009 Elsevier B.V. All rights reserved.

1. Introduction Films of chalcogenide glasses are currently a subject of systematic research, among other reasons, because of the changes in physical and chemical properties, which take place in samples after illumination or annealing [1–6]. These photo- and thermally-induced processes offer the possibility of using amorphous chalcogenide semiconductors for high-density information storage, high-resolution display devices and fabrication of diffractive optical elements [7]. In particular, As40 S60−x Sex ternary mixed glasses are attractive candidates for all these important technological applications [8–10]. The aim of the present paper is to study the changes in the structure and optical properties – the optical dispersion and absorption – of thermally-evaporated amorphous thin films of chemical composition As40 S20 Se40 , after illumination with bandgap light, in air, and after thermal annealing at a temperature near the glass-transition temperature, Tg . Structural properties have been deduced from both X-ray diffraction (XRD) patterns and Raman spectra of the thin-film chalcogenide samples. Optical properties have been derived from transmission spectra, at normal incidence, in the 400–2200 nm spectral range. The results obtained have been systematically compared with those found in the literature for amorphous layers of the

∗ Corresponding author. Tel.: +34 956 016318; fax: +34 956 016288. E-mail address: [email protected] (E. Márquez). 0254-0584/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2009.02.038

closely related binary stoichiometric compositions, As40 S60 (As2 S3 ) and As40 Se60 (As2 Se3 ) [2]. This investigation is of particular interest given the scarcity of literature on the interrelationship between the structure and optical constants of both untreated and treated ternary-chalcogenide-glass thin films. 2. Experimental details Amorphous As40 S20 Se40 thin films were prepared by vacuum thermal evaporation following the procedure described in detail in Ref. [11]. The deposition rate was ≈7 nm s−1 , as measured by the dynamic weighing method. Actual chemical composition of the thermally-evaporated chalcogenide films was found to be As39.1±1.9 S22.2±0.9 Se38.7±1.9 , on the basis of electron microprobe X-ray analysis, using a scanning electron microscope (JEOL, model JSM-820). Some virgin layers were illuminated in air, by a 500 W Hg arc lamp (Oriel, model 6285), through an infrared-cut filter, providing broadband ‘white’ light, with a high ultraviolet output, and with a light intensity of ≈40 mW cm−2 . It should be emphasized that around 3 h of illumination with the Hg lamp, was found to be enough to reach the saturation state, and that the temperature increase of the films during the irradiation period, as measured by a thermocouple affixed to the surface of the samples, was never more than ≈25 ◦ C. This particular observation gives unambiguous evidence of the athermal character of the observed photo-induced transformations. Other as-deposited layers were annealed at 160 ◦ C, for approximately 24 h, under a vacuum of about 10−3 Torr. The lack of crystallinity in the as-evaporated films was verified by XRD measurements (Philips, model PW-1820 diffractometer). After illumination or annealing, the XRD analysis was also performed in order to study the possible structural changes occurring during these two treatments. The Raman spectra were measured by using a Fourier transform IR spectrometer (Bruker, model IFS 55), with a Raman accessory (Bruker, model FRA 106); the resolution of the Raman spectrometer was 1 cm−1 . The optical transmission spectra were obtained at normal incidence by a double-beam UV/Vis/NIR spectrophotometer (PerkinElmer, model Lambda-19); the wavelength

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is composed, e.g., As4 S(Se)4 , As4 S(Se)3 and S(Se)n . These kinds of molecular fragments have also been reported in the case of ternary As–S–Se glassy alloys [20]. They obviously make difficult the cohesion between the structural layers, increasing the free volume in the material. On the other hand, these molecular species are unstable with respect to illumination or annealing. So, under light exposure or thermal annealing, their concentration is greatly reduced by polymerization and cross-linking with the amorphous network. Consequently, a stronger interaction between the structural layers, through an increased number of intermolecular chemical bonds involving neighbouring structural layers, is achieved – such bonds are of van der Waals type, between As atoms and S or Se chalcogen atoms. Thus, the structure and, therefore, the mass density of the treated chalcogenide film become almost identical to that of the corresponding bulk glass [12,13]. Lastly, it is important to note that in illuminated and annealed As40 S60 and As40 Se60 binary layers, a very strong change in the FSDP was reported by De Neufville et al. [2]. 3.2. Raman spectroscopy

Fig. 1. XRD patterns (CuK␣ radiation,  = 1.54 Å) of the untreated and treated amorphous As40 S20 Se40 films: (i) virgin; (ii) exposed to bandgap illumination, in air; and (iii) annealed in vacuum, at 160 ◦ C.

range analyzed was between 400 and 2200 nm. A surface-profiling stylus instrument (Sloan, model Dektak 3030) was also used to independently measure the film thickness, which was compared with the average thickness calculated from the transmission measurements. Furthermore, mass measurements were made by a microbalance (Mettler, model AE200) to check possible changes as a result of the illumination or annealing.

3. Results and discussion 3.1. X-ray diffraction Fig. 1 shows the XRD patterns corresponding to as-deposited, illuminated and annealed amorphous As40 S20 Se40 films. A significant feature of the diffraction data of covalently-bonded non-crystalline solids, which has been traditionally associated with the presence of medium-range order, is the first sharp diffraction peak (FSDP), occurring at values of the magnitude of the scattering vector, Q (=4sin/) = 1–2 Å−1 , depending on the material [12,13]. In the present case, this feature appears at 2 = 16.54◦ , or, equivalently, at Q = 1.17 Å−1 . Many attempts have been made to explain the origin of the FSDP in glasses. Elliott [14,15] has proposed an explanation in which the FSDP is ascribed to a chemical-order pre-peak (in the concentration–concentration partial structural factor, Scc (Q), in the Bhatia–Thornton formalism [16]), due to the interstitial volume around the cation-centred structural units. This association of the FSDP with correlations involving interstitial voids, plausibly explains the anomalous behaviour of this peak as a function of temperature and pressure. It can also be seen in Fig. 1 that after exposure of the virgin sample, the FSDP does disappear, while in the case of annealing, the intensity profile of the FSDP clearly decreases. These experimental results suggest, according to Elliott’s ideas, a diminution of the interstitial volume around the AsS3−n Sen (n = 0,1,2,3) pyramidal structural units, which form the layered network of the As40 S60−x Sex mixed glasses [17], i.e., a compaction of the amorphous structure (changes in the masses of the samples were not found after both treatments). It is well-known [1,12,18,19] that thermally-evaporated binary amorphous As–S and As–Se films often contain some of the molecular species of which the vapour

The Raman spectra of as-evaporated, exposed and annealed As40 S20 Se40 layers, are compared with the Raman spectrum of the bulk glass of the same chemical composition, in Fig. 2. A very good correlation between the positions of the main bands in these Raman spectra of bulk glass, virgin, illuminated and annealed films certainly exists. In this way, one can conclude that the same basic structural units form the glass matrix of a given composition. The main broad bands (220–260 and 320–380 cm−1 spectral regions) are, in turn, ‘structured’, that is, they are, in fact, the result of the overlapping of narrower bands. At the same time, it is observed that some additional weaker bands exist in the Raman spectrum of the virgin film, mainly, in 125–200 cm−1 spectral region. It should be pointed out that bulk glasses of the binary compositions As40 S60 and As40 Se60 show quite simple Raman spectra, with the main bands at 345 and 227 cm−1 , respectively, which are associated with vibrations of AsS3 and AsSe3 pyramidal structural units. As expected, the Raman spectrum corresponding to the bulk glass of the present ternary composition show a ‘two-phase structure’,

Fig. 2. Raman spectra of As40 S20 Se40 samples: (i) virgin; (ii) exposed to bandgap illumination, in air; (iii) annealed in vacuum, at 160 ◦ C; and (iv) bulk glass of the same chemical composition.

E. Márquez et al. / Materials Chemistry and Physics 115 (2009) 751–756

Fig. 3. Transmission as a function of photon energy, T(ω), of a wedge-shaped, as-evaporated As40 S20 Se40 chalcogenide film, with an average thickness of 769 ± 3 (0.4%) nm; T+ and T− are the top and bottom envelope curves, respectively. The film-on-substrate system is also shown in this figure.

i.e., two main bands with maxima at wavenumbers very close to the above-mentioned ones (see Fig. 2). Such a two-phase structure, is preserved in the case of the thin film. Nevertheless, a number of additional features can also be seen in the Raman spectrum corresponding to the untreated As40 S20 Se40 thin-film sample: the weak and narrow bands appearing in the 125–200 cm−1 spectral region, which are associated to the already indicated molecular species, As4 S(Se)4 , As4 S(Se)3 and S(Se)n [20,21]. Exposure or annealing of the virgin sample result in a very significant decrease in the intensity of all of these particular bands, which strongly supports the proposed idea of the photo- and thermally-induced polymerization of such molecular clusters. 3.3. Optical properties 3.3.1. Dispersion of the refractive index As-evaporated, illuminated and annealed amorphous As40 S20 Se40 films, have been geometrically and optically characterized from their corresponding optical transmission spectra, at normal incidence – Fig. 3 shows a typical optical transmission spectrum for an as-deposited film, along with its computer-drawn upper and lower envelopes [22]. Improved analytical expressions for the two envelopes of the transmission spectrum (for both uniform and non-uniform thickness thin films), which, importantly, takes into account the very often present (and not considered) weak absorption in the glass substrate [23], were introduced into the well-known envelope optical characterization method

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suggested by Swanepoel [24,25]. This new approach has allowed us to determine the average thickness of the wedge-shaped films ¯ and the refractive index, n, with accuracies better being studied, d, than 1%. It must be noted that, even though a ‘planetary rotation’ coating system was used, in order to reduce to some extent the non-uniformity in the thickness of the films, all the samples showed a ‘detectable’ lack of uniformity in thickness. Values of the thickness variation, d, at the extreme of the spectrophotometer light spot area (1 × 4 mm2 ), are listed in Table 1. All the details about the formulation and the algorithm used for the determination of the geometrical and optical parameters, can be found in our previous paper [23]. It has been calculated, using the optical characterization method mentioned above, that the exposure of the as-deposited sample to bandgap light induces a decrease in its average thickness of 2.1% (from 756 ± 2 (0.3%) nm down to 740 ± 3 (0.4%) nm), which is a sign of a photo-densification (photo-contraction) process. Similarly, the annealing of the virgin sample causes a decrease in its average thickness of 5.9% (from 769 ± 3 (0.4%) nm down to 724 ± 2 (0.3%) nm), because, as already indicated, illumination as well as annealing, takes the structure (and, therefore, the mass density) of the virgin film very close to the network structure of the corresponding bulk glass. It is stressed that in both the untreated and treated samples, the thickness determined by mechanical measurements with the stylus-based surface profiler, were in excellent agreement with the values of d¯ calculated by the optical method (the differences being always less than 2%). In Fig. 4, the final, more accurate values of the refractive index are plotted as a function of photon energy, ω. In this figure, it can be seen that n increases both upon illumination and upon annealing. An explanation that could account for this increase is presented below. On the other hand, the dispersion of the refractive index has been analyzed in terms of the Wemple–DiDomenico (WDD) model [26,27], which is based on the single-oscillator formula: n2 (¯hω) = 1 +

Eo Ed 2

Eo2 − (¯hω)

,

(1)

where Eo is the single-oscillator energy and Ed the dispersion energy or single-oscillator strength. By plotting (n2 − 1)−1 against (ω)2 and fitting a straight line, as shown in the inset of Fig. 4, Eo and Ed are directly determined from the slope, (Eo Ed )−1 , and the intercept, Eo /Ed , on the vertical axis. However, it must be pointed out that, due to the optical absorption, the experimental variation in the refractive index clearly departs from that given by Eq. (1), when opt the photon energy approaches the Tauc gap, Eg , which will be formally defined later, when the optical absorption edge is studied (see Fig. 4) [26,28]. The values of WDD dispersion parameters Eo and Ed , as well as the corresponding static  refractive index, n(0) (the refractive 1 + Ed /Eo ), for as-evaporated, exposed index at ω = 0, n(0) = and annealed films, are all listed in Table 1. The single-oscillator energy is an ‘average’ energy gap, and to a good approximation it opt opt scales with the optical bandgap Eg , Eo ≈ 2 × Eg , as was found by Tanaka [29] investigating well-annealed Asx S100−x chalcogenide glass layers. Moreover, an important achievement of the WDD model is that it relates the dispersion energy, Ed , to other physical

Table 1 Geometrical and optical parameters for the virgin (v) — virgin samples to be illuminated and to be annealed, respectively —, illuminated (i), and annealed (a), amorphous As40 S20 Se40 thin films studied. opt

State

d¯ (nm)

d (nm)

Eo (eV)

Ed (eV)

Nc

n(0)

Eg

v (→i)

756 ± 2 740 ± 3

8±1 17 ± 1

4.14 ± 0.01 3.91 ± 0.03

22.12 ± 0.06 22.92 ± 0.19

3.20 ± 0.30 3.32 ± 0.30

2.520 ± 0.001 2.620 ± 0.002

769 ± 3 724 ± 2

8±1 9±1

4.18 ± 0.02 3.88 ± 0.03

22.54 ± 0.13 23.35 ± 0.16

3.26 ± 0.30 3.38 ± 0.30

2.528 ± 0.001 2.649 ± 0.002

i

v (→a) a

opt

B1/2 (cm−1/2 eV−1/2 )

E04 (eV)

Eo /Eg

1.93 ± 0.01 1.88 ± 0.01

778 ± 1 825 ± 1

2.12 2.04

2.2 2.1

1.94 ± 0.01 1.89 ± 0.01

799 ± 1 841 ± 3

2.12 2.06

2.2 2.1

(eV)

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reported for low energy (IR-region) by De Neufville et al. [2], for the particular cases of binary As40 S60 and As40 Se60 films (+0.107 ± 0.01 and +0.092 ± 0.01, and +0.062 ± 0.01 and +0.058 ± 0.01, respectively). However, it is worth indicating that, although it has been found in both As40 S60 and As40 Se60 layers that the low-energy refractive index increases slightly more after exposure than after annealing, in the case of the static refractive index of As40 S20 Se40 layers, we have obtained just the opposite. The experimentally-obtained relative changes for n(0), +4.0% after illumination, and +4.8% after annealing, can be related to the increase in the mass density of the layer, , or, equivalently, to the decrease of its average thickness, according to the Lorentz–Lorenz equation [31,32], by the following expression: n(0) 6n2 (0) d¯ =− 2 , 2 n(0) (n (0) − 1)(n (0) + 2) d¯

Fig. 4. Refractive index as a function of photon energy, n(¯hω), for the virgin, exposed and annealed As40 S20 Se40 samples. The curves have been drawn from the Wemple–DiDomenico dispersion relationship. In the inset, the plot of the refractiveindex factor (n2 − 1)−1 vs. (ω)2 .

parameters, through a simple empirical relationship [26,27]: Ed = ˇNc Za Ne

(eV),

(2)

where Nc is the effective coordination number of the cation nearestneighbour to the anion, Za is the formal chemical valency of the anion, Ne is the effective number of valence electrons per anion, and ˇ is a two-valued constant, with either an ionic or a covalent value (ˇi = 0.26 ± 0.03 eV and ˇc = 0.37 ± 0.04 eV). As a consequence of both illumination and annealing, the single-oscillator strength, Ed , does increase. Considering that Ne = (40 × 5 + 60 × 6)/60 = 9 1/3, and Za = 2, and assuming that they should not be changed by either of the two treatments, it seems reasonable to ascribe this trend observed in the values of Ed , to an increase in the effective cation coordination number, Nc . It is well-known [13] that the structure of binary As2 Ch3 amorphous chalcogenides (Ch being a chalcogen atom) consists of locally twodimensional structural layers, formed by AsCh3 pyramidal units linked through a common two-fold coordinated chalcogen atom, and interacting with each other by weak intermolecular bonds. It has been shown [17,20,30] that this picture is also valid for the structure of ternary As40 S60−x Sex amorphous chalcogenides. Also, according to Wemple [27], interactions between structural layers through As atoms acting as ‘bonding points’ (forming As· · ·Ch intermolecular bonds), contribute to increase the As effective coordination number, and, thus, a higher-than-three value of Nc is expected. In particular, for As40 S60 bulk glass, Wemple suggests a value of Nc ≈ 3.2. On the other hand, the additional increase in Nc after both illumination and annealing, could be explained in terms of an increased van der Waals interaction between these structural layers, as a result of the polymerization and the higher structural compactness of the films, after both treatments. The effective values of Nc for the virgin, exposed and annealed samples, derived from Eq. (2), are all listed in Table 1. The changes found in the static refractive index upon illumination (n(0) = +0.100 ± 0.003), and upon annealing (n(0) = +0.121 ± 0.003), are, indeed, consistent with those

(3)

¯ d¯ = −/. The relative changes found in the average where d/ thickness, −2.1% and −5.9%, after illumination and after annealing, respectively, give place, according to Eq. (3), to relative changes in n(0) of +1.8% and +5.0%, respectively. There is an almost total agreement with the experimentally-obtained relative change in n(0) in the case of annealing, but, on the contrary, there is a notable difference in the case of exposure that we believe it is explained in terms of an increase in the number of heteropolar bonds, which are formed at the expense of homopolar bonds present in the various molecular species contained in the virgin film, giving rise finally to an increase in the effective polarizability of the amorphous material. 3.3.2. Absorption edge The dependence of the absorption coefficient, ˛, on photon energy is displayed in Fig. 5, using a convenient semi-logarithmic scale. As can be seen in this figure, the exposure to bandgap light causes a clear shift of the absorption edge of the as-deposited film to lower photon energies (i.e., photo-darkening). Similarly, the thermal annealing leads to a significant red shift of the absorp-

Fig. 5. Changes in the optical absorption edge of a virgin As40 S20 Se40 film, induced by either bandgap illumination or thermal annealing: a red shift takes place in both cases. In the inset, the determination of the Tauc gap by extrapolation in a (˛ω)1/2 vs. ω plot (Tauc extrapolation).

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tion edge of the virgin layer (i.e., thermal-darkening). In the strong-absorption region (where ˛ > ≈ 104 cm−1 ), which involves electronic transitions between the valence and conduction bands, the absorption coefficient is given according to the ‘non-direct electronic transition’ model proposed by Tauc [33], by the following quadratic equation: opt 2

˛(¯hω) = B

(¯hω − Eg ) h ¯ω

,

(4)

where B is a constant, which depends on the electronic transition opt probability, and Eg is the already introduced Tauc gap, being now opt

formally defined. The values of Eg and B1/2 can be readily derived from Eq. (4), by plotting (˛ω)1/2 vs. ω (Tauc plot). The excellent fit of the straight line corresponding to the (˛ω)1/2 vs. ω plot to the high-energy experimental points, shown in the inset of Fig. 5, indicates that the non-direct electronic transition is the mechanism responsible for the optical absorption in this high-energy spectral region, in the case of the as-deposited, and photo- and thermaldarkened As40 S20 Se40 layers. opt The values of the Tauc gap and slope, Eg and B1/2 , respectively, as well as the alternative optical gap, E04 , which represents the energy at which the absorption coefficient reaches the value of 104 cm−1 , for the glassy samples under study, are presented in Table 1. The change found in the optical gap E04 after annealing opt was E04 = −0.06 eV (Eg = −0.05 eV), while those reported by De Neufville et al. [2], E04 (As40 S60 ) and E04 (As40 Se60 ), were both −0.05 eV. In addition, the change in E04 upon illumination opt was E04 = −0.08 eV (Eg = −0.05 eV), and those found in Ref. [2] were E04 (As40 S60 ) = −0.07 eV and E04 (As40 Se60 ) = −0.06 eV. opt The decrease of both optical gaps, E04 and Eg , could be related to some extent to the lower average molar bond energy of the As–S heteropolar bonds, 260 kJ mol−1 [7], compared to that corresponding to the S–S homopolar, ‘wrong bonds’, 280 kJ mol−1 ; these S–S homopolar bonds are contained in some of the vapour molecular groups embedded in the atomic structure of the as-deposited sample. Furthermore, one can conclude from the results listed in Table 1, opt that the values of the WDD and Tauc gaps, Eo and Eg , respectively, satisfy fairly well the already introduced scaling relationship, opt Eo ≈ 2 × Eg . There is also a reasonably good agreement with the opt suggested by Tichá and Tichy´ [34], Eo = opt 1.25 × (1.2 + Eg ); in particular, in the case of the treated samples,

formula relating Eo and Eg

the agreement is nearly total. Last but not least, it is important to mention that the increase of the Tauc slope, derived from the values of B1/2 presented in Table 1, corresponding to exposure and annealing, +47 ± 2 cm−1/2 eV−1/2 and +42 ± 3 cm−1/2 eV−1/2 , respectively, are smaller than those reported by Tichy´ et al. [5] for the composition As2 S3 , +61 cm−1/2 eV−1/2 and +120 cm−1/2 eV−1/2 , respectively. This is interpreted as a larger decrease in the degree of structural disorder, as a result of both treatments, in the case of the As2 S3 glassy alloy [35] – the Tauc slope, B1/2 , is known to be inversely proportional to the degree of structural randomness of the amorphous network, which, as a consequence of the two treatments, is undoubtedly diminished. 3.3.3. Photo-oxidation process It must be pointed out that, as in the previously reported data on light-exposed As40 S60 and As40 Se60 chalcogenide layers [1,36], a phase separation, leading to the formation of a new phase of As2 O3 (arsenolite), on the surface of the As40 S20 Se40 films, has been found. Thus, the XRD pattern of the exposed samples (see Fig. 1) shows two Bragg peaks at 2 = 14.02◦ and 28.04◦ , respectively – their Miller indices and d-spacings are also indicated in Fig. 1 [37]. Regarding the crystal structure of arsenolite, it has also to be mentioned that the unit-cell dimension, a, in each direc-

Fig. 6. SEM images of the surface of a representative As–S–Se thin-film sample, after it was exposed to UV light, for about 3 h, at ambient conditions. A micro-crystal of As2 O3 (a typical octahedron of arsenolite [37]), formed as a result of a photo-oxidation effect, is displayed in this figure.

tion (cubic crystal system), for As2 O3 , is 11.073 Å, and its volume is V = a3 = 1357.8 Å3 . The observed surface effect, on the other hand, is closely related to the strong absorption for high photon energies, that the amorphous chalcogenide layers generally exhibit (see Fig. 5). Moreover, it should be taken into account that the spectral irradiance curve of the Hg arc lamp used in this work to illuminate the virgin samples, shows very high output peaks in the region of approximately 300–450 nm (corresponding to photon energies of about 4.1–2.8 eV) [38]. Therefore, a considerable part of the radiation emitted by the Hg lamp (mainly, the UV part), does not penetrate into the chalcogenide film. This is, indeed, the cause of the partial photo-degradation of the layer at its surface. As light exposure has been carried out in air, both photooxidation and photo-hydrolysis could take place (because of the presence of both oxygen and water vapour), which would result in the appearance of small, but clearly recognizable, arsenic trioxide micro-crystals [1,36]. So, the surface morphology of the illuminated samples was carefully examined by SEM, in order to identify these As2 O3 micro-crystals (see Fig. 6). Owing to the low concentration of them found on the surface of the samples, they should not significantly influence the present optical transmission measurements, nor, consequently, all the results derived from them. In addition, it has to be mentioned that no measurable compositional change was found in the light-exposed layer, by energy-dispersive spectroscopy. We finally propose that the observed photo-oxidation process of amorphous As40 S20 Se40 (As2 SSe2 ) films could be accounted for

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by considering the following light-induced chemical reaction: As2 SSe2 +

3 h ¯ω O2 −→As2 O3 + S + 2Se, 2

References (5)

taking place at the surface of the chalcogenide film. This photochemical reaction is absolutely similar to that previously suggested by De Neufville et al. [2], in the case of the photo-enhanced oxidation of the glassy composition As2 S3 , which can be schematically represented by: h ¯ω

(As2 S3 )amorphous + O2 −→(As2 O3 )crystalline + (S)crystalline.

(6)

This reaction (6) is thermodynamically favourable at room temperature (it is accelerated at higher temperatures), but it is essentially inhibited by some activation barrier in the absence of bandgap illumination. 4. Concluding remarks Thermally-evaporated amorphous As40 S20 Se40 thin films undergo structural transformations when exposed to bandgap illumination, with a Hg arc lamp, in air (i.e., a photo-structural transformation), and when annealed at a temperature near Tg (i.e., a thermo-structural transformation). The disappearance, on the one hand, and the decrease, on the other hand, of the FSDP in the XRD diffraction pattern, upon illumination and upon annealing, respectively, has been interpreted as a diminution of the interstitial volume around the AsS3−n Sen (n = 0,1,2,3) pyramidal structural units, through the polymerization of the various molecular species embedded in the atomic structure of the as-evaporated chalcogenide films, leading finally to a more cross-linked amorphous network. This structural change is certainly consistent with the decrease of the average thickness found for the samples, whether illuminated or annealed. Also, clear indications of a photo-oxidation process, on the surface of the light-exposed films, has been found: arsenic trioxide micro-crystals were unequivocally identified, visually, by SEM, and, structurally, by XRD measurements. Changes in the optical properties – optical dispersion and absorption – of the virgin chalcogenide films after bandgap illumination as well as after annealing, have been systematically studied in this work. In particular, the refractive index of the as-evaporated layer increases with either of the two treatments. Another relevant finding is the noticeable red shift of the optical absorption edge of the as-deposited sample, as a result of both light exposure (that is, a photo-darkening process has taken place), and thermal annealing (in other words, a thermal-darkening effect). Acknowledgments This work has been supported by the MCYT (Spain) and FEDER (EU), under FIS2005-01409 research project.

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