Hydrogen is the element of choice to passivate dangling bonds in amorphous and crystalline semiconductors, like Si, Ge and SiGe because of its single-electron atomic structure. By hydrogenation better electro-optical characteristics are obtained. Hydrogenated a-Si, a-Ge and a-SiGe are still deeply investigated [1-4] because of their promising employement in many devices, like solar cells, IR detectors. However, the behavior of H is not fully understood and predictable when the above materials are submitted to illumination or thermal treatment [5, 6]. Though a huge literature exists on a-Si and a-SiGe alloy, a-Ge and a-Si/a-Ge multilayers have been little studied. The latter ones can be candidate to form the a-SiGe alloy by annealing them so as to intermix Si and Ge and create the alloy [7, 8]. This work deals with the effect of annealing on the evolution of the structure and associated changes in the H configuration in hydrogenated amorphous a-Si/a-Ge multilayers deposited by RF sputtering. Results on annealed hydrogenated single layers of a-Si and a-Ge are also presented as they turned out to be useful to interpret the H behavior in the multilayers. The hydrogenated a-Si/a-Ge multilayers (MLs) were deposited by RF sputtering, from high purity crystalline Si and Ge targets on (100) Si substrate, with a mixture of high purity Ar and H gases. The MLs consisted of fifty couples of alternating a-Si and a-Ge layers, 3 nm thick each, yielding a ML thickness of 300 nm. Hydrogenation was achieved by letting H flow into the sputtering chamber during the whole deposition period. H flow rates were 0.4, 0.8 and 1.5 ml/min. Annealing was done in high purity (99.999%) argon at 350 °C or 400 °C for 1, 4 and 10 hours. Single layers of a-Si and a-Ge, 40 nm thick, were also grown, under the same conditions, to measure the H incorporation efficiency. The samples were analysed by AFM, ERDA (Elastic Recoil Detection Analysis), Infrared (IR) Absorption, SEM, TEM and STEM-HAADF. For ERDA the 1.6 MeV 4He+ beam available at the 5 MeV Van de Graaf accelerator of Budapest was applied to measure the H in the samples. The incorporated H content in the unannealed samples increased with increasing H flow rate with an asymptotic behavior at the highest H flow rates, reaching a value of ~17 at % and ~ 7 at % for a-Si and a-Ge, respectively, at the H flow rate of 1.5 ml/min. The H was mostly incorporated as Si-H and Si-H2 in the Si layers and as Ge-H in the Ge layers. The not hydrogenated MLs did not show any structural modification upon annealing. However, structural modifications in the shape of blistering were observed in the annealed hydrogenated MLs, with formation of surface bubbles by plastic deformation whose density and size increased with increasing H content, for the same annealing conditions [9, 10]. Along with such structural degradation a change of the H bonding configuration to Si and Ge also occurred. In fact, the Si-H, Si-H2 and Ge-H peaks in the IR absorbance spectra disappeared upon annealing, to an extent dependent on the annealing conditions, indicating that H was released free to the lattice and that de-passivation of the dangling bonds occurred. The complete break of the H bonds to the Ge atoms took place at annealing times shorter than for Si, confirming that the binding energy of the Ge-H bond is lower than that of the Si-H bond. This suggests that the observed bubbles and blistering are due to local accumulation of the liberated H. Surface blistering due to bubbles was also observed in the single a-Si and a-Ge layers. For the highest H content it was observed that in the a-Ge layers the great majority of the bubbles have transformed into craters, i.e. they blew up because of a high internal gas pressure, while in a-Si they underwent such transformation only to a very much less extent. In the annealed a-Ge samples containing craters ERDA showed that the H remained in the layers was only 15% of the incorporated one whilst it was about 65% in a-Si which suggests that H had escaped though the craters. This further confirms that the bubbles have formed by accumulation of H. In a-Ge hydrogen release and formation of H bubbles is thus more efficient and occurs at an earlier time, with consequent earlier bubble explosion, than in a-Si. The anticipated rupture of the Ge layer has also to be partially ascribed to the lower mechanical strength of the Ge lattice with respect to Si. Therefore, blistering of the a-Si/a-Ge ML nanostructures very likely primarily started by accumulation of H liberated from the Ge atoms. Nucleation sites for the growth of the bubbles are expected to be nanocavities present in the amorphous phase. Calculations of the density of H2 molecules contained in the bubbles by the Wan et al. model of lenticular crack [11] show that it is 9.9 and 5.7 % of the H remained in a-Si layers and a-Si/a-Ge MLs, respectively, after annealing. [1] G. G. Pethuraja, R. E. Welser, A. K. Sood, C. Lee, N.J. Alexander, H. Efstathiadis, P. Haldar, J. L. Harvey, Mater. Scie. Appl. 3 (2012) 67 [2] J. Müllerová, L. Prusáková, M. Netrvalová, V. Vavrunková, P. Sutta, Appl. Surf. Sci. 256 (2010) 5667 [3] Y. Hamakawa, J. Non-Cryst. Solids 352 (2006) 863 [4] K. W. Jobson, J.-P. R. Wells, R. E. I. Schropp. D. A. Carder, P. J. Philips, J. I. Dijkhuis, Phys. Rev. B 73 (2006) 155202 [5] J. D. Cohen, Solar Energy Mater. Solar Cells 78 (2003) 399 [6] P. Agarwal and S. C. Agarwal, Phil. Mag. B 80 (2000) 1327 [7] T. Sameshima, H. Watanabe, H. Kanno, T. Sadoh, M. Miyao, Thin Solid Films 487 (2005) 67 [8] M. S. Abo Ghazala, Physica B 293 (2000) 132 [9] C. Frigeri, L. Nasi, M. Serényi, A. Csik, Z. Erdélyi, D.L. Beke, Superlatt. Microstruct. 45 (2009) 475 [10] C. Frigeri, M. Serényi, N. Q. Khánh, A. Csik, Z. Erdélyi, L. Nasi, D. L. Beke, H.-G. Boyen, Nanoscale Research Letters 6 (2011) 189 [11] K.-T. Wan, R. G. Horn, S. Courmont, B. R. Lawn, J. Mater. Res. 8 (1993) 1126
Evolution of the structure and hydrogen configuration in annealed hydrogenated a-Si/a-Ge multilayers and layers
2012
Abstract
Hydrogen is the element of choice to passivate dangling bonds in amorphous and crystalline semiconductors, like Si, Ge and SiGe because of its single-electron atomic structure. By hydrogenation better electro-optical characteristics are obtained. Hydrogenated a-Si, a-Ge and a-SiGe are still deeply investigated [1-4] because of their promising employement in many devices, like solar cells, IR detectors. However, the behavior of H is not fully understood and predictable when the above materials are submitted to illumination or thermal treatment [5, 6]. Though a huge literature exists on a-Si and a-SiGe alloy, a-Ge and a-Si/a-Ge multilayers have been little studied. The latter ones can be candidate to form the a-SiGe alloy by annealing them so as to intermix Si and Ge and create the alloy [7, 8]. This work deals with the effect of annealing on the evolution of the structure and associated changes in the H configuration in hydrogenated amorphous a-Si/a-Ge multilayers deposited by RF sputtering. Results on annealed hydrogenated single layers of a-Si and a-Ge are also presented as they turned out to be useful to interpret the H behavior in the multilayers. The hydrogenated a-Si/a-Ge multilayers (MLs) were deposited by RF sputtering, from high purity crystalline Si and Ge targets on (100) Si substrate, with a mixture of high purity Ar and H gases. The MLs consisted of fifty couples of alternating a-Si and a-Ge layers, 3 nm thick each, yielding a ML thickness of 300 nm. Hydrogenation was achieved by letting H flow into the sputtering chamber during the whole deposition period. H flow rates were 0.4, 0.8 and 1.5 ml/min. Annealing was done in high purity (99.999%) argon at 350 °C or 400 °C for 1, 4 and 10 hours. Single layers of a-Si and a-Ge, 40 nm thick, were also grown, under the same conditions, to measure the H incorporation efficiency. The samples were analysed by AFM, ERDA (Elastic Recoil Detection Analysis), Infrared (IR) Absorption, SEM, TEM and STEM-HAADF. For ERDA the 1.6 MeV 4He+ beam available at the 5 MeV Van de Graaf accelerator of Budapest was applied to measure the H in the samples. The incorporated H content in the unannealed samples increased with increasing H flow rate with an asymptotic behavior at the highest H flow rates, reaching a value of ~17 at % and ~ 7 at % for a-Si and a-Ge, respectively, at the H flow rate of 1.5 ml/min. The H was mostly incorporated as Si-H and Si-H2 in the Si layers and as Ge-H in the Ge layers. The not hydrogenated MLs did not show any structural modification upon annealing. However, structural modifications in the shape of blistering were observed in the annealed hydrogenated MLs, with formation of surface bubbles by plastic deformation whose density and size increased with increasing H content, for the same annealing conditions [9, 10]. Along with such structural degradation a change of the H bonding configuration to Si and Ge also occurred. In fact, the Si-H, Si-H2 and Ge-H peaks in the IR absorbance spectra disappeared upon annealing, to an extent dependent on the annealing conditions, indicating that H was released free to the lattice and that de-passivation of the dangling bonds occurred. The complete break of the H bonds to the Ge atoms took place at annealing times shorter than for Si, confirming that the binding energy of the Ge-H bond is lower than that of the Si-H bond. This suggests that the observed bubbles and blistering are due to local accumulation of the liberated H. Surface blistering due to bubbles was also observed in the single a-Si and a-Ge layers. For the highest H content it was observed that in the a-Ge layers the great majority of the bubbles have transformed into craters, i.e. they blew up because of a high internal gas pressure, while in a-Si they underwent such transformation only to a very much less extent. In the annealed a-Ge samples containing craters ERDA showed that the H remained in the layers was only 15% of the incorporated one whilst it was about 65% in a-Si which suggests that H had escaped though the craters. This further confirms that the bubbles have formed by accumulation of H. In a-Ge hydrogen release and formation of H bubbles is thus more efficient and occurs at an earlier time, with consequent earlier bubble explosion, than in a-Si. The anticipated rupture of the Ge layer has also to be partially ascribed to the lower mechanical strength of the Ge lattice with respect to Si. Therefore, blistering of the a-Si/a-Ge ML nanostructures very likely primarily started by accumulation of H liberated from the Ge atoms. Nucleation sites for the growth of the bubbles are expected to be nanocavities present in the amorphous phase. Calculations of the density of H2 molecules contained in the bubbles by the Wan et al. model of lenticular crack [11] show that it is 9.9 and 5.7 % of the H remained in a-Si layers and a-Si/a-Ge MLs, respectively, after annealing. [1] G. G. Pethuraja, R. E. Welser, A. K. Sood, C. Lee, N.J. Alexander, H. Efstathiadis, P. Haldar, J. L. Harvey, Mater. Scie. Appl. 3 (2012) 67 [2] J. Müllerová, L. Prusáková, M. Netrvalová, V. Vavrunková, P. Sutta, Appl. Surf. Sci. 256 (2010) 5667 [3] Y. Hamakawa, J. Non-Cryst. Solids 352 (2006) 863 [4] K. W. Jobson, J.-P. R. Wells, R. E. I. Schropp. D. A. Carder, P. J. Philips, J. I. Dijkhuis, Phys. Rev. B 73 (2006) 155202 [5] J. D. Cohen, Solar Energy Mater. Solar Cells 78 (2003) 399 [6] P. Agarwal and S. C. Agarwal, Phil. Mag. B 80 (2000) 1327 [7] T. Sameshima, H. Watanabe, H. Kanno, T. Sadoh, M. Miyao, Thin Solid Films 487 (2005) 67 [8] M. S. Abo Ghazala, Physica B 293 (2000) 132 [9] C. Frigeri, L. Nasi, M. Serényi, A. Csik, Z. Erdélyi, D.L. Beke, Superlatt. Microstruct. 45 (2009) 475 [10] C. Frigeri, M. Serényi, N. Q. Khánh, A. Csik, Z. Erdélyi, L. Nasi, D. L. Beke, H.-G. Boyen, Nanoscale Research Letters 6 (2011) 189 [11] K.-T. Wan, R. G. Horn, S. Courmont, B. R. Lawn, J. Mater. Res. 8 (1993) 1126I documenti in IRIS sono protetti da copyright e tutti i diritti sono riservati, salvo diversa indicazione.
