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Influence of laser fluence in ArF-excimer laser assisted crystallisation of a-SiGe:H films
Applied Surface Science 208±209 (2003) 358±363
In¯uence of laser ¯uence in ArF-excimer laser assisted crystallisation of a-SiGe:H ®lms S. Chiussia,*, E. LoÂpeza, J. Serraa,c, P. GonzaÂleza, C. Serraa, B. LeoÂna, F. Fabbrib, L. Fornarinib, S. Martellib,1 a
Dpto. FõÂsica Aplicada, Universidade de Vigo, Lagoas Marcosende, E-36200 Vigo, Spain b ENEA Frascati, Via Enrico Fermi 27, I-00044 Frascati (Rome), Italy c CACTI Universidade de Vigo, Lagoas Marcosende, E-36200 Vigo, Spain
Abstract Polycrystalline silicon germanium (poly-SiGe) coatings are drawing increasing attention as active layers in solar cells, bolometers and various microelectronic devices. As a consequence, alternative low-cost production techniques, capable to produce such alloys with uniform and controlled grain size, become more and more attractive. Excimer laser assisted crystallisation, already assessed in thin ®lm transistor production, has proved to be a valuable ``low-thermal budget'' technique for the crystallisation of amorphous silicon. Main advantages are the high process quality and reproducibility as well as the possibility of tailoring the grain size in both, small selected regions and large areas. The feasibility of this technique for producing poly-SiGe ®lms has been studied irradiating hydrogenated amorphous SiGe ®lms with spatially uniform ArF-laser pulses of different ¯uences. Surface morphology, structure and chemical composition have been extensively characterised, demonstrating the need of using a ``step-by-step'' process and a careful adjustment of both, total number of shots and laser ¯uence at each ``step'' in order to diminish segregation effects and severe damages of the ®lm surface and of segregation effects. # 2002 Elsevier Science B.V. All rights reserved. Keywords: a-SiGe:H; Poly-SiGe; ArF-excimer laser; Laser-CVD; Excimer laser crystallisation
1. Introduction Polycrystalline silicon germanium (poly-SiGe) alloys have proved to be promising active materials for room temperature microbolometers, solar cells and various microelectronic devices [1±3]. The main advantages of poly-SiGe compared to poly-Si are a lower thermal conductivity that improves the thermal *
Corresponding author. Tel.: 34-98-681-2216; fax: 34-98-681-2201. E-mail address: [email protected] (S. Chiussi). 1 Present address: Centro Sviluppo Materiali, Via di Castel Romano 100, I-00128 Roma, Italy.
insulation of the device, a high temperature coef®cient of resistance and the possibility of tailoring the band gap according to the requirements of the desired optoand microelectronic devices. These material characteristics and the fact that SiGe is easy to micro-machine, are expected to drastically increase their performance and to permit their down-scaling, maintaining the compatibility with the IC silicon technology. Until now, poly-SiGe ®lms have been mainly produced through conventional chemical vapour deposition (CVD) processes [1] that require relatively high substrate temperatures (at least 650 8C) or solid phase crystallisation (SPC) of amorphous SiGe coatings (a-SiGe). This crystallisation technique needs slightly
0169-4332/02/$ ± see front matter # 2002 Elsevier Science B.V. All rights reserved. doi:10.1016/S0169-4332(02)01399-5
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
lower substrate temperatures but high processing times of at least several hours for crystallising a-SiGe on SiO2 substrates [4]. The development of fast alternative processes with lower thermal budget for enabling the use of low-cost substrates like glass and simplifying the integration of such ®lms in CMOS technology is therefore strongly motivated. Among such alternative processes, those using laser-assisted techniques like laser induced chemical vapour deposition (LCVD) and laser assisted crystallisation are the most attractive ones. ArF-LCVD in parallel con®guration for example has proved to be a feasible ``soft'' deposition technique for producing uniform amorphous materials with acceptable growth rates at low substrate temperatures [5±8]. The subsequent crystallisation of such amorphous ®lms throughexcimer laser radiation is an ultra-rapid annealing process and has already been used for producing polycrystalline silicon coatings on glass as well as partially relaxed heteroepitaxial SiGe ®lms on Si(1 0 0) [9±12]. Excimer laser crystallisation (ELC) in particular was already demonstrated to be economically effective and has been established for the production of poly-Si based solar cells and thin ®lm transistors (TFTs) [9]. This technique has also been successfully simulated through theoretical models for predicting the optimum process parameters leading to uniform and adjustable grain sizes [10]. However, due to segregation effects and stoichiometry dependent optical and thermodynamic properties, the experience based on poly-Si is not readily applicable to poly-SiGe. Moreover, despite the efforts on investigating the heteroepitaxial growth of SiGe-®lms on Si(1 0 0) [11], very little is known about the crystallisation processes of these alloys on amorphous substrates. A previously performed ®rst feasibility study of an ``all laser assisted'' process using ArF-LCVD for growing hydrogenated amorphous SiGe coating (a-SiGe:H) on glass and ArF-ELC for crystallising them, has shown that poly-SiGe can be produced [8]. In this process, a ``step-by-step'' crystallisation procedure with increasing laser ¯uence in each step has been used in order to subsequently dehydrogenise and crystallise the aSiGe:H ®lm for preventing explosive H-effusion. This work presents a detailed study on the role of the laser ¯uence in this crystallisation process, con®rming that such a ``step-by-step'' procedure is
359
indispensable for the crystallisation of a-SiGe:H ®lms, but also demonstrating that this procedure promotes segregation of Ge and roughening of the surface. 2. Experimental Poly-SiGe ®lms have been achieved irradiating 500 nm thick a-SiGe:H ®lms, grown on corning (7059) glass substrates through ArF-LCVD in parallel con®guration (substrate above the beam). Both processes have been carried out in the same conventional HV-system (base pressure 50 mPa) and using the same ArF-excimer laser (Lambda Physik 220i) [6,8]. During the LCVD process, the substrate temperature has been kept at 320 8C and the total pressure at 0.53 kPa. The 193 nm radiation inducing the reactions in the gas phase (1 sccm Si2H6 and 0.8 sccm GeH4 in 650 sccm He) has been adjusted to a power density of 0.5 W/cm2. After exposing the ®lms to air for producing a thin native oxide cap layer, the coatings have been irradiated in an inert gas atmosphere (600 sccm He at 0.53 kPa) with intensity homogenised 193 nm pulses (Exitech ``¯y-eyes''-homogenizer EXHS-700D, spot size 2:8 mm 5:6 mm, pulse length 20 ns). The irradiation has been carried out at room temperature using different laser ¯uences according to a ``step-bystep'' process starting with a laser ¯uence of 50 mJ/ cm2 and ending at ®nal values of 100, 200, 280, and 440 mJ/cm2. Composition of the ®lms has been determined by FT-IR-spectroscopy (MB100-Bomem), Xray photoelectron spectroscopy (XPS, ESCALAB 250iXL-VG Scienti®c), and time of ¯ight-secondary ion mass spectrometry (TOF-SIMS, TRIFT III-physical electronics). Surface morphology and roughness have been analysed by Pro®lometry (Dektak3ST-Veeco), scanning electron microscopy (SEM, XL30-PHILIPS) and AFM (atomic force microscopy, discoverer-Topometrix). The ®lm crystallinity has been studied by Raman spectroscopy (RFS100-Bruker), XRD (X-ray diffraction; Geiger¯ex-Rigaku), and transmission electron microscopy (TEM: CM20-PHILIPS). 3. Results and discussion From a-Si:H crystallisation it is known that the irradiation of H-rich ®lms with ¯uences of E >
360
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
Fig. 1. (a) SEM image showing the formation of randomly distributed holes in the coating, (b) dark ®eld TEM image with corresponding electron-diffraction pattern (in-set) indicating the partial crystallisation to ®ne grained polycrystalline material.
150 mJ/cm2 leads to an explosive H-effusion damaging the ®lm surface [13,14]. Similar results have now been obtained after irradiating 500 nm thick a-SiGe:H ®lms with ¯uences of E > 100 mJ/cm2. As it can be observed in the SEM image (Fig. 1a), the coating is completely perforated. The formation of the randomly distributed holes in the coating can easily be explained through the explosive effusion of H from the originally very smooth (rms 4 nm) ``as-deposited'' a-SiGe:H ®lms produced through ArF-LCVD in parallel con®guration [8]. On the other hand, TEM analysis (Fig. 1b) shows also a partial crystallisation to ®ne grained polycrystalline material as it can also be seen from the corresponding electron-diffraction pattern (in-set of Fig. 1b). It can be assumed that the ultrarapid crystallisation and the associated increase of density of the coating during the explosive hydrogen effusion might also enhance the formation of holes in the ®lm. The fact, that these effects appear at lower ¯uences than for irradiating pure a-Si:H is consistent with the more than 100 K lower effusion temperature of hydrogen bonded to Ge than hydrogen bonded to Si. [8,13]. An approach to solve this problem is to apply a ``step-by step'' process with gradually increasing ¯uence. According to previous experimental results of FT-IR-analysis and numerical simulation [8], ¯uences below the melting threshold of a surface layer initiates a partial dehydrogenation of the coating. Above this melting threshold, which is around
60±80 mJ/cm2, the increasing ¯uence produces deeper molten pools that are accompanied by a longer duration of the molten state and a heating of the underlying layers thus inducing their dehydrogenation. However, it is expected that the pulsed laser assisted crystallisation does not totally dehydrogenate the ®lms and that the residual hydrogen passivates grain boundary defects, thus avoiding the use of typical ``post-laser treatments'' with hydrogen [12,14]. The detailed study of the ®lm composition after different stages of the process has been achieved through XPS depth-pro®le analyses. Fig. 2 shows the evolution of the Si2p and the Ge3d peak of a representative sample indicating that both Si and Ge seem to have lost their homogeneous distribution in depth. Representing the concentration of both elements versus depth (Fig. 3) reveals the formation of a twolayer system. An ``upper''-layer with increasing Ge concentration towards the surface of the coating and an uniform ``underlying``-layer with the Ge concentration of the original a-SiGe:H ®lm. It can also be observed that the thickness of the ``upper''-layer increases with the ®nal ¯uence and that the gradient increases with the number of pulses. The one-dimensional numerical simulation of the ELC process [9] correlated the thickness of the ``upper''-layer with the depth of the molten pool, corroborating that the segregation is forced by the repetitive melt/solidi®cation cycles of the ``step-by-step'' process.
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
361
Fig. 2. Depth-evolution of the Si2p and Ge3d peaks of a sample irradiated in six steps (50, 75, 100, 125, 160, 200 mJ/cm2) with 10 pulses per step.
The prediction that the substrate surface remains practically thermally unaffected for the used range of ¯uences [8] has been corroborated by the absence of additional oxygen or barium that, in case of overheating the substrate, should have diffused from the glass into the coating. AFM analysis of the surface showed that the roughness increased with the ®nal ¯uence H®n from rms 4 nm for Hfin 100 mJ/cm2 and 6 nm for Hfin 200 mJ/cm2 to 40 nm for Hfin 280 mJ/cm2 indicating the beginning of island formation. The severe damage of the surface and the strong segregation of
Ge towards the surface can be seen more clearly if excessive number of pulses and high ¯uences are used. As it can be seen in the following SEM and AFM images (Fig. 4), a sample irradiated in 17 steps up to a ®nal ¯uence of 440 mJ/cm2 shows formation of islands on the 170 nm rough coating. This island formation should be caused by the fact that higher ¯uences produces an overheating of the ®lm, thus slowing down the solidi®cation velocity of the alloy (<4 m/s) so that the partitioning of Ge is no longer negligible [15]. The regions with higher Ge concentration cannot solidify until the Ge atoms segregates
Fig. 3. XPS depth-pro®le analysis of samples at different stages of the process. Samples have been irradiated (a) in 3 steps (50, 75, 100 mJ/ cm2) with 10 pulses per step, (b) in 6 steps (50, 75, 100, 125, 160, 200 mJ/cm2) with 10 pulses per step, (c) in 12 steps (50, 75, 100, 105, 130, 150, 165, 185, 220, 250, 270, 280 mJ/cm2) with 1 pulses per step.
362
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
Fig. 4. (a) SEM and (b) AFM image of a sample irradiated with an excessive number of pulses and elevated ®nale ¯uence (Hfin 440 mJ/ cm2).
causing preferential Ge aggregates on the surface. TOF-SIMS line scan analysis of the substrate surface ®nally con®rmed this fact revealing that the islands observed in the SEM image are enriched in Germanium.
the molten pool. Finally, the irradiation of the coatings with elevated ¯uences and number of pulses clearly shows the formation of Ge rich islands on the surface of the sample and, consequently, strong roughening of the surface.
4. Conclusion
Acknowledgements
Hole formation due to explosive hydrogen effusion during crystallising a-SiGe:H coatings on glass with ¯uences >100 mJ/cm that heating up the sample surface above the melting threshold, con®rmed the need of a ``step-by-step'' process a-SiGe:H. The study of the ¯uence dependence in such a process revealed that several melt/crystallisation cycles have to be performed in order to obtain an effective dehydrogenation and that a two-layer system consisting in an ``upper''layer with a Ge gradient and an ``underlying'' layer with the original SiGe alloy is formed. The results, consistent with the numerical simulation indicate that at ¯uences <80 mJ/cm2 only dehydrogenation takes place and that above this threshold a crystallisation of a thin surface layer occurs. An increasing ¯uence causes further dehydrogenation from deeper layers and the subsequent crystallisation of the H-poor ``upper''-layers with the segregation of Ge towards the surface. The thickness of these ``upper''-layers corresponds to the depth of the molten pool, indicating the Ge segregation is caused by the solidi®cation of
This work was partially supported by EU as well as by Spanish contracts MAT2000-1050, XUGA32107BB92DOG211, UV62903I5F4, PGIDT01PX130301PN, PR405A2001/35-0. The authors wish to thank U. Kosch (FH-O/O/W-Emden), J.B. RodrõÂguez (CACTI, Univ. Vigo), T. Sulima (BW-Univ. MuÈnchen) and A. Abalde for their extensive technical help and helpful discussions. TOF-SIMS measurements have been performed by Scott Bryan (Physical Instruments Minneapolis).
References [1] S. Sedky, P. Fiorini, M. Caymax, A. Verbist, C. Baert, Sens. Actuators A66 (1998) 193. [2] S. Sedky, P. Fiorini, M. Caymax, C. Baert, L. Hermans, R. Mertens, IEEE Elect. Device Lett. 19 (1998) 376. [3] N. Yamauchi, R. Reif, J. Appl. Phys. 75 (1994) 3235. [4] I. Tsunoda, T. Nagata, A. Kenjo, T. Sadoh, M. Miyao, Mater. Sci. Eng. B89 (2002) 336. [5] I.P. Herman, Chem. Rev. 89 (1989) 1323.
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363 [6] S. Chiussi, P. GonzaÂlez, J. Serra, B. LeoÂn, M. PeÂrez-Amor, Appl. Surf. Sci. 106 (1996) 75. [7] E. LoÂpez, S. Chiussi, J. Serra, P. GonzaÂlez, C. Serra, B. LeoÂn, M. PeÂrez-Amor, Appl. Surf. Science, in press. [8] S. Chiussi, C. Serra, J. Serra, P. GonzaÂlez, B. LeoÂn, S. Urban, G. AndraÈ, J. Bergmann, F. Falk, F. Fabbri, L. Fornarini, S. Martelli, F. Rinaldi, Appl. Surf. Sci. 186 (2002) 166. [9] V.V. Gupta, H.J. Song, J.S. Im, Appl. Phys. Lett. 71 (1997) 99. [10] S.U. Campisano, D.C. Jacobson, J.M. Poate, A.G. Cullis, N.G. Chew, Appl. Phys. Lett. 46 (1985) 846.
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[11] N. Frangis, J. VanLanduyt, R. Larciprete, S. Martelli, E. Borsella, S. Chiussi, J. Castro, B. LeoÂn, Appl. Phys. Lett. 72 (1998) 2877. [12] P. Lengsfeld, N.H. Nickel, W. Fuhs, Appl. Phys. Lett. 76 (2000) 1680. [13] K. Suzuki, M. Takahashi, M. Saitoh, Appl. Phys. A69 (1999) S263. [14] P. Mei, J.B. Boyce, M. Hack, R.A. Lujan, R.I. Johnson, G.B. Anderson, D.K. Fork, S.E. Ready, Appl. Phys. Lett. 64 (1994) 1132. [15] Q. Yu, P. Clancy, J. Cryst. Growth 149 (1995) 45.
In¯uence of laser ¯uence in ArF-excimer laser assisted crystallisation of a-SiGe:H ®lms S. Chiussia,*, E. LoÂpeza, J. Serraa,c, P. GonzaÂleza, C. Serraa, B. LeoÂna, F. Fabbrib, L. Fornarinib, S. Martellib,1 a
Dpto. FõÂsica Aplicada, Universidade de Vigo, Lagoas Marcosende, E-36200 Vigo, Spain b ENEA Frascati, Via Enrico Fermi 27, I-00044 Frascati (Rome), Italy c CACTI Universidade de Vigo, Lagoas Marcosende, E-36200 Vigo, Spain
Abstract Polycrystalline silicon germanium (poly-SiGe) coatings are drawing increasing attention as active layers in solar cells, bolometers and various microelectronic devices. As a consequence, alternative low-cost production techniques, capable to produce such alloys with uniform and controlled grain size, become more and more attractive. Excimer laser assisted crystallisation, already assessed in thin ®lm transistor production, has proved to be a valuable ``low-thermal budget'' technique for the crystallisation of amorphous silicon. Main advantages are the high process quality and reproducibility as well as the possibility of tailoring the grain size in both, small selected regions and large areas. The feasibility of this technique for producing poly-SiGe ®lms has been studied irradiating hydrogenated amorphous SiGe ®lms with spatially uniform ArF-laser pulses of different ¯uences. Surface morphology, structure and chemical composition have been extensively characterised, demonstrating the need of using a ``step-by-step'' process and a careful adjustment of both, total number of shots and laser ¯uence at each ``step'' in order to diminish segregation effects and severe damages of the ®lm surface and of segregation effects. # 2002 Elsevier Science B.V. All rights reserved. Keywords: a-SiGe:H; Poly-SiGe; ArF-excimer laser; Laser-CVD; Excimer laser crystallisation
1. Introduction Polycrystalline silicon germanium (poly-SiGe) alloys have proved to be promising active materials for room temperature microbolometers, solar cells and various microelectronic devices [1±3]. The main advantages of poly-SiGe compared to poly-Si are a lower thermal conductivity that improves the thermal *
Corresponding author. Tel.: 34-98-681-2216; fax: 34-98-681-2201. E-mail address: [email protected] (S. Chiussi). 1 Present address: Centro Sviluppo Materiali, Via di Castel Romano 100, I-00128 Roma, Italy.
insulation of the device, a high temperature coef®cient of resistance and the possibility of tailoring the band gap according to the requirements of the desired optoand microelectronic devices. These material characteristics and the fact that SiGe is easy to micro-machine, are expected to drastically increase their performance and to permit their down-scaling, maintaining the compatibility with the IC silicon technology. Until now, poly-SiGe ®lms have been mainly produced through conventional chemical vapour deposition (CVD) processes [1] that require relatively high substrate temperatures (at least 650 8C) or solid phase crystallisation (SPC) of amorphous SiGe coatings (a-SiGe). This crystallisation technique needs slightly
0169-4332/02/$ ± see front matter # 2002 Elsevier Science B.V. All rights reserved. doi:10.1016/S0169-4332(02)01399-5
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
lower substrate temperatures but high processing times of at least several hours for crystallising a-SiGe on SiO2 substrates [4]. The development of fast alternative processes with lower thermal budget for enabling the use of low-cost substrates like glass and simplifying the integration of such ®lms in CMOS technology is therefore strongly motivated. Among such alternative processes, those using laser-assisted techniques like laser induced chemical vapour deposition (LCVD) and laser assisted crystallisation are the most attractive ones. ArF-LCVD in parallel con®guration for example has proved to be a feasible ``soft'' deposition technique for producing uniform amorphous materials with acceptable growth rates at low substrate temperatures [5±8]. The subsequent crystallisation of such amorphous ®lms throughexcimer laser radiation is an ultra-rapid annealing process and has already been used for producing polycrystalline silicon coatings on glass as well as partially relaxed heteroepitaxial SiGe ®lms on Si(1 0 0) [9±12]. Excimer laser crystallisation (ELC) in particular was already demonstrated to be economically effective and has been established for the production of poly-Si based solar cells and thin ®lm transistors (TFTs) [9]. This technique has also been successfully simulated through theoretical models for predicting the optimum process parameters leading to uniform and adjustable grain sizes [10]. However, due to segregation effects and stoichiometry dependent optical and thermodynamic properties, the experience based on poly-Si is not readily applicable to poly-SiGe. Moreover, despite the efforts on investigating the heteroepitaxial growth of SiGe-®lms on Si(1 0 0) [11], very little is known about the crystallisation processes of these alloys on amorphous substrates. A previously performed ®rst feasibility study of an ``all laser assisted'' process using ArF-LCVD for growing hydrogenated amorphous SiGe coating (a-SiGe:H) on glass and ArF-ELC for crystallising them, has shown that poly-SiGe can be produced [8]. In this process, a ``step-by-step'' crystallisation procedure with increasing laser ¯uence in each step has been used in order to subsequently dehydrogenise and crystallise the aSiGe:H ®lm for preventing explosive H-effusion. This work presents a detailed study on the role of the laser ¯uence in this crystallisation process, con®rming that such a ``step-by-step'' procedure is
359
indispensable for the crystallisation of a-SiGe:H ®lms, but also demonstrating that this procedure promotes segregation of Ge and roughening of the surface. 2. Experimental Poly-SiGe ®lms have been achieved irradiating 500 nm thick a-SiGe:H ®lms, grown on corning (7059) glass substrates through ArF-LCVD in parallel con®guration (substrate above the beam). Both processes have been carried out in the same conventional HV-system (base pressure 50 mPa) and using the same ArF-excimer laser (Lambda Physik 220i) [6,8]. During the LCVD process, the substrate temperature has been kept at 320 8C and the total pressure at 0.53 kPa. The 193 nm radiation inducing the reactions in the gas phase (1 sccm Si2H6 and 0.8 sccm GeH4 in 650 sccm He) has been adjusted to a power density of 0.5 W/cm2. After exposing the ®lms to air for producing a thin native oxide cap layer, the coatings have been irradiated in an inert gas atmosphere (600 sccm He at 0.53 kPa) with intensity homogenised 193 nm pulses (Exitech ``¯y-eyes''-homogenizer EXHS-700D, spot size 2:8 mm 5:6 mm, pulse length 20 ns). The irradiation has been carried out at room temperature using different laser ¯uences according to a ``step-bystep'' process starting with a laser ¯uence of 50 mJ/ cm2 and ending at ®nal values of 100, 200, 280, and 440 mJ/cm2. Composition of the ®lms has been determined by FT-IR-spectroscopy (MB100-Bomem), Xray photoelectron spectroscopy (XPS, ESCALAB 250iXL-VG Scienti®c), and time of ¯ight-secondary ion mass spectrometry (TOF-SIMS, TRIFT III-physical electronics). Surface morphology and roughness have been analysed by Pro®lometry (Dektak3ST-Veeco), scanning electron microscopy (SEM, XL30-PHILIPS) and AFM (atomic force microscopy, discoverer-Topometrix). The ®lm crystallinity has been studied by Raman spectroscopy (RFS100-Bruker), XRD (X-ray diffraction; Geiger¯ex-Rigaku), and transmission electron microscopy (TEM: CM20-PHILIPS). 3. Results and discussion From a-Si:H crystallisation it is known that the irradiation of H-rich ®lms with ¯uences of E >
360
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
Fig. 1. (a) SEM image showing the formation of randomly distributed holes in the coating, (b) dark ®eld TEM image with corresponding electron-diffraction pattern (in-set) indicating the partial crystallisation to ®ne grained polycrystalline material.
150 mJ/cm2 leads to an explosive H-effusion damaging the ®lm surface [13,14]. Similar results have now been obtained after irradiating 500 nm thick a-SiGe:H ®lms with ¯uences of E > 100 mJ/cm2. As it can be observed in the SEM image (Fig. 1a), the coating is completely perforated. The formation of the randomly distributed holes in the coating can easily be explained through the explosive effusion of H from the originally very smooth (rms 4 nm) ``as-deposited'' a-SiGe:H ®lms produced through ArF-LCVD in parallel con®guration [8]. On the other hand, TEM analysis (Fig. 1b) shows also a partial crystallisation to ®ne grained polycrystalline material as it can also be seen from the corresponding electron-diffraction pattern (in-set of Fig. 1b). It can be assumed that the ultrarapid crystallisation and the associated increase of density of the coating during the explosive hydrogen effusion might also enhance the formation of holes in the ®lm. The fact, that these effects appear at lower ¯uences than for irradiating pure a-Si:H is consistent with the more than 100 K lower effusion temperature of hydrogen bonded to Ge than hydrogen bonded to Si. [8,13]. An approach to solve this problem is to apply a ``step-by step'' process with gradually increasing ¯uence. According to previous experimental results of FT-IR-analysis and numerical simulation [8], ¯uences below the melting threshold of a surface layer initiates a partial dehydrogenation of the coating. Above this melting threshold, which is around
60±80 mJ/cm2, the increasing ¯uence produces deeper molten pools that are accompanied by a longer duration of the molten state and a heating of the underlying layers thus inducing their dehydrogenation. However, it is expected that the pulsed laser assisted crystallisation does not totally dehydrogenate the ®lms and that the residual hydrogen passivates grain boundary defects, thus avoiding the use of typical ``post-laser treatments'' with hydrogen [12,14]. The detailed study of the ®lm composition after different stages of the process has been achieved through XPS depth-pro®le analyses. Fig. 2 shows the evolution of the Si2p and the Ge3d peak of a representative sample indicating that both Si and Ge seem to have lost their homogeneous distribution in depth. Representing the concentration of both elements versus depth (Fig. 3) reveals the formation of a twolayer system. An ``upper''-layer with increasing Ge concentration towards the surface of the coating and an uniform ``underlying``-layer with the Ge concentration of the original a-SiGe:H ®lm. It can also be observed that the thickness of the ``upper''-layer increases with the ®nal ¯uence and that the gradient increases with the number of pulses. The one-dimensional numerical simulation of the ELC process [9] correlated the thickness of the ``upper''-layer with the depth of the molten pool, corroborating that the segregation is forced by the repetitive melt/solidi®cation cycles of the ``step-by-step'' process.
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
361
Fig. 2. Depth-evolution of the Si2p and Ge3d peaks of a sample irradiated in six steps (50, 75, 100, 125, 160, 200 mJ/cm2) with 10 pulses per step.
The prediction that the substrate surface remains practically thermally unaffected for the used range of ¯uences [8] has been corroborated by the absence of additional oxygen or barium that, in case of overheating the substrate, should have diffused from the glass into the coating. AFM analysis of the surface showed that the roughness increased with the ®nal ¯uence H®n from rms 4 nm for Hfin 100 mJ/cm2 and 6 nm for Hfin 200 mJ/cm2 to 40 nm for Hfin 280 mJ/cm2 indicating the beginning of island formation. The severe damage of the surface and the strong segregation of
Ge towards the surface can be seen more clearly if excessive number of pulses and high ¯uences are used. As it can be seen in the following SEM and AFM images (Fig. 4), a sample irradiated in 17 steps up to a ®nal ¯uence of 440 mJ/cm2 shows formation of islands on the 170 nm rough coating. This island formation should be caused by the fact that higher ¯uences produces an overheating of the ®lm, thus slowing down the solidi®cation velocity of the alloy (<4 m/s) so that the partitioning of Ge is no longer negligible [15]. The regions with higher Ge concentration cannot solidify until the Ge atoms segregates
Fig. 3. XPS depth-pro®le analysis of samples at different stages of the process. Samples have been irradiated (a) in 3 steps (50, 75, 100 mJ/ cm2) with 10 pulses per step, (b) in 6 steps (50, 75, 100, 125, 160, 200 mJ/cm2) with 10 pulses per step, (c) in 12 steps (50, 75, 100, 105, 130, 150, 165, 185, 220, 250, 270, 280 mJ/cm2) with 1 pulses per step.
362
S. Chiussi et al. / Applied Surface Science 208±209 (2003) 358±363
Fig. 4. (a) SEM and (b) AFM image of a sample irradiated with an excessive number of pulses and elevated ®nale ¯uence (Hfin 440 mJ/ cm2).
causing preferential Ge aggregates on the surface. TOF-SIMS line scan analysis of the substrate surface ®nally con®rmed this fact revealing that the islands observed in the SEM image are enriched in Germanium.
the molten pool. Finally, the irradiation of the coatings with elevated ¯uences and number of pulses clearly shows the formation of Ge rich islands on the surface of the sample and, consequently, strong roughening of the surface.
4. Conclusion
Acknowledgements
Hole formation due to explosive hydrogen effusion during crystallising a-SiGe:H coatings on glass with ¯uences >100 mJ/cm that heating up the sample surface above the melting threshold, con®rmed the need of a ``step-by-step'' process a-SiGe:H. The study of the ¯uence dependence in such a process revealed that several melt/crystallisation cycles have to be performed in order to obtain an effective dehydrogenation and that a two-layer system consisting in an ``upper''layer with a Ge gradient and an ``underlying'' layer with the original SiGe alloy is formed. The results, consistent with the numerical simulation indicate that at ¯uences <80 mJ/cm2 only dehydrogenation takes place and that above this threshold a crystallisation of a thin surface layer occurs. An increasing ¯uence causes further dehydrogenation from deeper layers and the subsequent crystallisation of the H-poor ``upper''-layers with the segregation of Ge towards the surface. The thickness of these ``upper''-layers corresponds to the depth of the molten pool, indicating the Ge segregation is caused by the solidi®cation of
This work was partially supported by EU as well as by Spanish contracts MAT2000-1050, XUGA32107BB92DOG211, UV62903I5F4, PGIDT01PX130301PN, PR405A2001/35-0. The authors wish to thank U. Kosch (FH-O/O/W-Emden), J.B. RodrõÂguez (CACTI, Univ. Vigo), T. Sulima (BW-Univ. MuÈnchen) and A. Abalde for their extensive technical help and helpful discussions. TOF-SIMS measurements have been performed by Scott Bryan (Physical Instruments Minneapolis).
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