High Quality a-ge:h Films and Devices Through Enhanced Plasma Chemistry

Transcription

1 Mater. Res. Soc. Symp. Proc. Vol Materials Research Society 0989-A04-04 High Quality a-ge:h Films and Devices Through Enhanced Plasma Chemistry Erik V. Johnson, and Pere Roca i Cabarrocas LPICM (CNRS, UMR 7647), Ecole Polytechnique, Palaiseau Cedex, F-91128, France ABSTRACT We present a material study on RF PECVD-grown a-ge:h showing that thin films of such material can be produced without using the conventional techniques of high power density or powered-electrode substrate placement. We demonstrate the production of material with PDS signatures superior to material produced at ten times higher power density. This is achieved through the use of Ar and H 2 dilution and by growing the films at high pressures under conditions where nanoparticles and nanocrystals formed in the gas phase contribute significantly to the growth as confirmed by HRTEM. The conditions described result in material which demonstrates activated conduction down to room temperature. Additionally, the quality of the material has been demonstrated through its application in n-i-p diodes. A spectral response at 0.9um of 0.38 and an AM1.5 efficiency of 2.1% have been demonstrated utilizing an absorber layer thickness of only 60nm. INTRODUCTION The goal of producing PECVD-grown, device-quality a-ge:h for use as a low-bandgap semiconductor has lead to important discoveries about the ideal growth conditions of the material. It is generally accepted that the best quality a-ge:h is grown under conditions that are in stark contrast with those ideal for a-si:h, a similar tetrahedrally-bonded amorphous semiconductor. For example, the quality of a-ge:h is greatly improved when it is grown under conditions of high H 2 dilution of germane (GeH 4 ) and under high power conditions, whereas the growth of a-si:h under these conditions (without an accompanying increase in pressure) leads to low quality, porous material. High quality a-ge:h has been produced through deposition on the powered cathode of asymmetric RF-PECVD systems [1,2] and in ECR systems [3-6]. Recently, hot-wire CVD using undiluted GeH 4 has been employed to grow high-quality films [7], demonstrating that high H 2 -dilution is not a strict requirement but that the conditions for good quality films are more complicated. The causes behind the higher quality of films grown under the above-listed conditions are still under dispute. Authors variously ascribe the high quality of the material to be due to 1) copious amounts of atomic hydrogen arriving at the growth surface, 2) ion bombardment during growth, 3) growth in the gas depletion regime for germane, and/or (4) heated precursors arriving at the surface. In this work, we promote this growth environment through the use of simultaneous Ar+H 2 dilution of GeH 4 and an elevated growth pressure. The metastable argon ions (Ar*) created through electron collision will ionize GeH 4 molecules, increase the density of atomic H, and release energy to the surface through relaxation of the metastable state. As well, under the high-pressure conditions employed, plasma-formed and heated nanoparticles will contribute to the growth, resulting in ìpolymorphousî (pm-) material. A low substrate

2 temperature is also employed to encourage the incorporation of H to passivate dangling bonds. The use of conditions promoting polymorphous material has been previously shown to enhance the stability of the transport properties of pm-sige:h films under light-soaking [8]. It has been demonstrated in an accompanying study [9] that pm-ge:h grown in conditions promoting nanoparticle incorporation into the film can be applied as a 60nm thick i- layer in n-i-p photodiodes demonstrating 2.1% efficiency at AM1.5 and spectral response in the near infrared (900nm) of 0.38 despite illumination through the n-layer. We investigate this device-quality pm-ge:h through optical and electronic measurements to further examine the nature of the material and to explain the origin of these high quality films produced in nonstandard conditions. EXPERIMENT The films were deposited in the ARCAM reactor [10], a multiplasma, monochamber RF- PECVD system operating at MHz. The walls of the chamber were heated to the same temperature as the substrate holder. The substrates were placed on the un-powered, grounded electrode, in contrast with results from the literature showing that the best films are produced through placement on the powered electrode. This substrate holder can be rotated in and out of the plasma/deposition zone. Two pressure series of films on Corning Glass Type 1737 (CG1737) and two films on carbon-covered TEM grids were deposited under the conditions listed in Table 1. Table 1 Growth Conditions Material Series A Material Series B HRTEM A HRTEM B Argon Flow Rate (FR Argon ) 200 sccm 200 sccm 200 sccm 200 sccm Hydrogen Flow Rate (FR H2 ) 200 sccm 200 sccm 200 sccm 200 sccm Germane Flow Rate (FR GeH4 ) 6 sccm 6 sccm 6 sccm 5 sccm Substrate/Chamber Temp. 175 C 175 C 175 C 175 C Pressure (mtorr) Interelectrode Distance (mm) Power Density (mw/cm 2 ) The material is characterized through HRTEM, Raman scattering, temperature-dependent conductivity, Spectroscopic Ellipsometry (SE) and photothermal deflection spectroscopy (PDS). RESULTS AND DISCUSSION HRTEM In Fig. 1 are presented images of two films grown under different growth conditions. The first two images are of Sample HRTEM A (conditions listed in Table 1), grown at the lower power density and pressure. This sample was exposed to air after deposition. Apparent in Fig.1(a) is an agglomerate of larger nanoparticles 10-20nm in diameter (possibly collected during rotation out of the deposition zone) surrounded by amorphous particles sized from 2-4 nm. The greater magnification image in Fig.1 (b) shows these contrasted regions. These regions have been previously interpreted as island-like growth [11]. However, cross-sectional HRTEM results (not

3 included in this study) suggest that they are the result of plasma-formed nanoparticles being deposited onto the growth surface and incorporated into the film. Fig. 1(c) shows the image from Sample HRTEM B, grown at higher power and pressure and with a slightly lower germane flow (5 vs 6 sccm). Clearly visible are the nanocrystalline domains, up to 10nm in diameter, despite the films being on the order of 10nm thick. Figure 1 - HRTEM images of two pm/nc-ge:h samples. Subfigures (a) and (b) show Sample A, consisting of amorphous nanoparticles 2-4 nm in diameter, and subfigures (c) and (d) show Sample B, consisting of nanocrystals up to 10nm in size. See text for discussion. Temperature Dependent Conductivity The temperature dependence of the dark conductivity was measured for the films grown in Material Series A and B. The samples were measured in vacuum, and the curves presented in Fig. 2 are measurements taken after the samples had been put through one temperature cycle (heating to 120 C, cooling to 40 C). The results for the samples grown at 60mW/cm 2 are plotted in Fig. 2(a), and these results can be fit well to activated conduction, indicating transport predominantly through extended states, at temperatures descending down to 40 C. The results for the samples grown at 120mW/cm 2 show a similar behaviour, with the exception of the sample grown at the highest pressures (2000 and 2600 mtorr). These samples show two

4 temperature regions of conduction: a lower temperature region with smaller activation energy than the rest of the samples, and a high temperature region with larger activation energy. (a) (b) Figure 2 ñ Temperature dependent conductivity data for samples grown at (a) 60mW/cm 2 and (b) 120mW/cm 2 This behaviour coincides with an elevated nanocrystalline content, as determined by Raman scattering and by ex-situ spectroscopic ellipsometry. However, the samples grown at low pressure and low power - which show a small crystalline signal from Raman ñ nevertheless exhibit activated conduction throughout the temperature range. In general, the samples grown at 60mW/cm 2 exhibit higher activation energies than those grown at 120mW/cm 2. The activation energies for the 60mW/cm 2 samples, however, are smaller than half the gap. This result is consistent with the fact that intrinsic a-ge:h is slightly n-type, and additionally, since no efforts were made to cap the material, some post-deposition O-contamination may have occurred. The nanoparticle-based growth confirmed by HRTEM emphasizes the large surface area available for post-deposition oxidation. The sample grown at 60mW/cm 2 and at 1250 mtorr exhibits an activation energy of 0.3eV, the largest amongst these samples. PDS and SE The mobility gap density of states was measured using Photothermal Deflection Spectroscopy (PDS), and the absorption coefficient spectra derived from this measurement are presented in Fig. 3. The clearest trend can be noted from the absorption coefficient value at 0.9eV, shown in Fig. 4a for the samples from Material Series A. In general, samples that show a microcrystalline character from Raman exhibit the highest absorption coefficients at 0.9eV. This result is consistent with the plasma-formation and incorporation of nanocrystals into the film, and these showing a remnant of the smaller c-ge gap. The films were also characterized through ex-situ spectroscopic ellipsometry (SE). A decrease in the maximum value of <ε2> - the imaginary part of the pseudo-dielectric function - with increasing pressure is noted for both sets of samples, as shown in Fig. 4(b). Comparison with results for high-quality a-ge:h from other groups [7,12,13] indicate that the material grown at low power and low pressure is comparably dense, but at increasing power is less so.

5 (a) (b) Figure 3 ñ Absorption spectrum from PDS for samples grown at (a) 60mW/cm 2 and (b) 120mW/cm 2. The open squares indicate points extracted from Ref. [1] for samples grown at 80mW/cm 2 (upper square) and 800mW/cm 2 (lower square). (a) (b) Figure 4 ñ (a) Absorption at 0.9eV as measured by PDS and approximate volume fraction of crystalline Ge (fractional area under Raman peak at 300cm -1 ) for sample grown at 60 mw/cm 2, and (b) maximum value of <ε 2 > from the acquired SE curve. Figure 5 ñ J-V Characteristic for pm-ge:h cell grown under elevated plasma chemistry conditions Use of the material grown at 60mW/cm 2 and at lower pressure (1 Torr) in photodiodes has resulted in an elevated infrared photovoltaic response and a very high short circuit current of J SC = 20.6mA/cm 2 despite an extremely thin i-layer (60nm) [8]. A J-V characteristic from such a device is presented in Fig. 5.These results are comparable to those achieved by growing diodes with ECR [14] and by mounting the substrate on the asymmetric, powered electrode [15].

6 CONCLUSIONS This material study elucidates the promising results from the incorporation of this material in devices, as presented in Ref. [9]. Of the a-ge:h grown for this study, only ìdevice-quality conditionsî material exhibited - activated conduction from 40 C to 120 C (indicating a single transport mechanism) - high material density (from SE) - a small but existing nanocrystalline fraction (from TEM, SE, Raman, and PDS) These properties coincided with the use of plasma conditions that encourage the formation and heating of nanoparticles in the plasma, as well as their incorporation into the film. It is proposed that under these conditions, the heating of Ge nanoparticles in the plasma encourages the formation of strong Ge-Ge bonds, while the low substrate temperature allows the incorporation of H in the film. The strong-bond formation role is typically played by ion-bombardment or by hot atoms reforming weak bonds at the surface, as proposed by Doyle et al [16] and as experimentally supported by results from hot-wire growth [6]. The results of this study show that it can also be enacted through enhancing plasma chemistry by simultaneous Ar+H 2 dilution. ACKNOWLEDGMENTS EVJ acknowledges the generous support of NSERC. REFERENCES [1] F.H. Karg, H. Bohm, and K. Pierz, J. Non-Cryst. Solids 114, 477 (1989). [2] W.A. Turner, et al, J. Appl. Phys. 67, 7430 (1990). [3] T. Aoki, S. Kato, Y. Nishikawa, and M. Hirose, J. Non-Cryst. Solids 114, 798 (1989). [4] A. Matsuda and K. Tanaka J. Non-Cryst. Solids 97 & 98, 1367 (1987). [5] N. Shibata, A. Tanabe, J. Hanna, S. Oda, and I. Shimizu, Jpn. J. Appl. Phys. 25, L540 (1986). [6] Xuejin Niu, V.L. Dalal, J. Appl. Phys. 98, (2005). [7] L. Zanzig, W. Beyer, and H. Wagner Appl. Phys. Lett. 67, 1567 (1995). [8] M.E. Guenier, J.P. Kleider, P. Chatterjee, P. Roca i Cabarrocas, and Y. Poissant, J. Appl. Phys. 92, 4959 (2002). [9] E.V. Johnson and P. Roca i Cabarrocas, (2007) Solar Energy Mater. and Solar Cells, doi: /j.solmat [10] P.Roca i Cabarrocas, J.B. ChÈvrier, J. Huc, A. Lioret, J.Y. Parey, and J.P.M. Schmitt, J. Vac. Sci. Technol. A 9, 2331 (1991). [11] B. Drevillon and C. Godet, J. Appl. Phys. 64, 145 (1988). [12] J.R. Blanco, P.J. McMarr, K. Vedam, and R.C. Ross, J. Appl. Phys. 60, 3724 (1986). [13] J.E. Yehoda, B. Yang, K. Vedam, and R. Messier, J. Vac. Sci. Technol. A 6, 1631 (1988). [14] J. Zhu, V.L. Dalal, M.A. Ring, and J.J. Guitterez, J.D. Cohen, J. Non-Cryst. Solids , 651 (2004). [15] W. Kusian, E. Gunzel and R.D. Plattner, Solar Energy Materials 23, 303 (1991). [16] J.R. Doyle, D.A. Doughty, and A. Gallagher, J. Appl. Phys. 69, 4169 (1991).