a-SiN0.72: H superlattices prepared by ion beam assisted reactive deposition

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Journal of Non-Crystalline Solids 155 (1993) 221-230 North-Holland

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a-SiNo.s8 • H/a-SiNo.72" H superlattices prepared by ion beam

assisted reactive deposition S. Zhang, Z.Y. W a n g a n d D.E. Brodie ( GWP) 2, University of Waterloo, Waterloo, ON N2L 3G1, Canada Received 15 July 1992 Revised manuscript received 12 November 1992

Superlattices have been fabricated with barrier layers of a-SiN0.72 : H and well layers of a-SiN05s : H. Si was evaporated from an e-gun source onto glass substrates and an a m m o n i a ion beam from a K a u f m a n type ion source simultaneously bombarded the substrate, introducing both N and H into the a-Si layers. The N concentration can be changed quickly by varying the ion beam current with all other deposition conditions held constant. The superlanice structure was observed directly either by transmission electron microscopy in cross-section, or by scanning electron microscopy. Also, low angle X-ray diffraction observations indicate that a superlattice structure has been formed. The transmission electron microscopy pictures show that layers have been deposited. The optical gaps, measured for the resulting devices, depend on the thickness of the well layer as well as the barrier width when these layers are thin. The optical gap increases with decreasing well widths when well widths are less than about 3 n m with fixed 6 nm thick barrier layers. On the other hand, the optical gap decreases when the barrier width is reduced to less than 5 nm and the well width is fixed at 1.5 nm. A q u a n t u m size effect does not appear to account for these two observations; however, they can be explained by non abrupt well-barrier layer interfaces. Both the dark conductivity and the photoconductivity depend on the well and barrier widths and the superlattices display photoconductivity decay and persistent photoconductivity.

1. Introduction Since Abeles and Tiedje [1] made the first amorphous semiconductor superlattice, many groups have investigated the properties of these superlattices with different structures and fabrication methods. For example, superlattices aS i : H / a - S i N x :H have been reported by Ibaraki and Fritzsche [2], Beaudoin et al. [3] and Miyazaki et al. [4], of a - S i : H / a - S i G % :H by Tiedje et al. [5] and of a-Si:H/a-SiCx :H by Bernhard et al. [6]. Glow discharge has been the most widely used method for preparing these superlattices, but photo-CVD has also been used [7]. There are arguments on whether the quantum size effect is present in amorphous semiconductor superlattices. When the well width is < 5 nm, most Correspondence to: Dr D.E. Brodie, D e p a r t m e n t of Physics, University of Waterloo, Waterloo, O N N2L 3G1, Canada.Tel: + 1-519 885 1211. Telefax: + 1-519 746 8115.

authors observed a blue shift in the optical gaps (e.g., Abeles et al. [1] and Miyazaki et al. [4]). This shift is also observed in the photoluminescence spectrum [8]. The increasing optical gap with decreasing well width was attributed to a quantum size effect observed in crystalline semiconductor superlattices [1,4,8]. Compelling evidence for this assumption is obtained from the results of thermal modulation absorption spectra [9] and resonant tunnelling phenomena [10]. However, this assumption has been challenged by other authors. Collins and Huang [11] suggested from theoretical considerations that a blue shift can occur as an artifact of the Tauc plot. This prediction was tested in experiments by Beaudoin et al. [3] and they observed that the optical gap only increased when the barrier width was < 1.5 nm, if a Cody plot @(a/hu) vs. hu) was used. They speculated that this optical increase was caused by the presence of interface states. Bernhard et al. [6] reported that a blue shift can be

0022-3093/93/$06.00 © 1993 - Elsevier Science Publishers B.V. All rights reserved

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S. Zhang et al. / a-SiN058 : n / a-SiNo.72 : H superlattices

obtained from coherent multilayer thin film optics and a quantum size effect was not needed. Their analysis was able to account for the changes in the optical gap measured using Tauc plots. Also, the critical well width of = 5 nm observed for the onset of the blue shift was questioned by Raikh et al. [12]. Their calculations suggest that quantization effects should only be present when the well width becomes less than = 2 nm, i.e., twice the carrier mean-free-path which is often assumed to be = 1 nm. This number is smaller than that generally observed in the experiments. Miyazaki and co-workers [10,13] interpreted their observations as due to resonant tunnelling through the thin potential barriers. However, Jiang and Hwang [14] showed that quantized energy levels in the well layers did not play a role in the carrier tunnelling observed in their six-layer structures. In these publications, work has focused on the effects of the well width and the role of the barrier width in superlattices has received less attention. In a crystalline semiconductor superlattice, the barrier layer width is not expected to affect the change in the optical gap for a fixed well layer thickness, as long as the barrier width does not become smaller than the critical tunnelling thickness or larger than the critical thickness required by the lattice mismatch. However, in a superlattice made with amorphous semiconductors, the barrier width could be an important parameter which may provide information concerning the origin of the blue shift in the optical absorption [6]. In this paper, we report a new technique for fabricating amorphous semiconductor superlattices, a-SiN0.58 : H/a-SiN0.72 : H multilayer structures were made using ion-beam-assisted reactive deposition. The effect of well layer width, as well as barrier layer width, on some of the electrical and optical properties of these structures were investigated.

2. Experimental Both the well layers (SiNo.ss:H) and barrier layers (a-SiNo.72 : H) were deposited by evaporat-

ing the silicon from an electron beam source while the growing film was bombarded with low energy ammonia ions from a Kaufman-type, broad-beam ion source. A hot filament in front of the source neutralized the beam to provide minimum beam spreading and substrate charging. The superlattice structures were deposited on Corning 7059 glass substrates. Single layers of the two different compositions were also deposited on polished c-Si wafers mounted on one corner of the glass, for IR absorption studies. For all depositions, substrate temperatures were fixed at 168°C and the ion beam voltage, discharge voltage and the accelerating voltage for the ion source were kept at 50, 120 and 420 V, respectively. The silicon deposition rate was monitored with a quartz crystal microbalance shielded from the ion source and was maintained at 4 n m / m i n . Under these conditions, the resulting nitrogen concentration, N N, in the layers is determined primarily by the magnitude of the ion beam current. The value of N N and hence the optical gap, was switched by changing the ion beam current and leaving all other deposition parameters fixed. The total thicknesses of all of the superlattices fabricated for this study were near 200 nm. Infrared absorption spectra from the single layer films deposited on the c-Si substrates were used to estimate the nitrogen and hydrogen concentrations, using the formula

N=Afa/,o d,o,

(1)

where a is the absorption coefficient calculated from the IR transmission spectrum, o~ is the wavenumber and A is a constant. For the N concentration, we used A = 6.3 x 1018 and the integration extended over the 840 cm -1 peak [15]. For the hydrogen concentration bound to Si, we used A = 1.4 x 1020 and the integration was over the 2170 cm-1 peak [15]. The number of H atoms bound to N was calculated using A = 2.8 × 1020 and the peak [16] at 3340 cm -1. These concentrations, determined from IR measurements, have not been checked by another independent method and we estimate an error of 20% in them due mainly to the difficulty in choosing the correct baseline.

S. Zhang et aL / a-SiNo.58 : H / a-SiNo. 72 : H superlattices

For a complete superlattice deposition, Si was evaporated without interruption. In order to achieve the sharpest possible transition when switching from one layer to the next, the following sequence of events (requiring about 10 s) was introduced: the shutter was closed, the ion beam current was changed and stabilized, and the shutter was opened again. A cross-sectional sample of the superlattice structure was prepared for T E M observations as follows: a diamond saw is used to cut the superlattice sample into 250 ~m strips and these strips are mounted in a film-to-film configuration; an ultrasonic cutter is used to prepare 3 mm diameter cylinders, suitable for a T E M sample holder, and slices from this cylinder are polished to a thickness of 10 Ixm using diamond paste; a dimple is introduced by mechanical polishing to give the sample a concave shape; finally, the specimen is milled by a focused Ar ion beam until the sample is transparent to the electron beam of the TEM. Samples for SEM observations were made by etching holes with low angle walls, using a mask and a low energy Ar ion beam, then overcoating this with a 20 nm layer of gold. Conductivity parallel to the plane of the layers was measured by preparing planar samples as follows: the conductivity sample was defined by a mask which protected a rectangular strip of the superlattice structure from an Ar ion beam which etched the sides of the strip; then, A1 for the electrical contacts was vacuum deposited along the two edges. The integrated well layer thickness was used to calculate the conductivity from the measured current. Errors in the conductivity measurements for the planar samples are = 10% and this arises mainly from the determination of the sample thickness in each case. However, samples with low room temperature dark conductivities ( < 10 -12 S / c m ) will have larger uncertainties, because these conductivities were determined by extrapolation from higher temperature dark conductivities. Photoconductivity and the photoconductivity degradation were measured using a mercury light source (i.e., an eraser source for an erasable, programmable, read only memory device) which gave a light intensity at the sample of = 60

223

txW/cm 2. Electrical activation energies were obtained from least square fits of the data to a straight line in plots of log (conductivity) vs. 1/T and the error in these values is _+0.05 eV. Optical gaps were obtained from Tauc plots (i.e., ~ vs. hu), although the applicability of a Cody plot (i.e., ~/[a/hu] vs. hu) was tested. The reproducibility of the optical gap determinations for different samples prepared in the same way was within +0.1 eV.

3. Results

3.1. An individual layer Figure 1 is a typical IR transmission spectrum for a single layer of a-SiN x : H. Except for differences in the magnitude of the absorption peaks obtained from films used for barriers or wells, all spectra obtained resembled those in fig. 1. Following the assignment given by Bustarret et al. [17], the peak at 3340 cm -1 is due to the N - H stretching mode and the peak at 2170 c m - 1 is the H - S i stretching mode modified by back-bonded nitrogen. The shoulder at 1150 cm -1 is due to the N - H wagging and rocking modes. The 840 and 450 cm-1 peaks are, respectively, the asymmetric and symmetric N - S i stretching modes. Us-

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Fig. I. A n infrared transmission spectrum for one of the a-SiN x :H single layer films. This spectrum shows the absorption peaks that are observed in the spectra for either the well or barrier layer materials which m a k e up the superlatticc structurcs. The peak assignments arc given in the text.

224

S. Zhang et al. / a-SiNo.58 : H ~ a-SiNo. 72 : H superlattices

ing eq. (1) and the peak at 840 cm -1 with a silicon concentration [15] of 5 × 1022/cm 3, the ratio of nitrogen to silicon is estimated to be 0.72 for the barrier layers and 0.58 for the well layers. The hydrogen concentration bound to Si is about 30 and 35 at.% for the barrier layers and well layers, respectively. In addition, a few at.% H is bonded to N. Using optical absorption and Tauc plots, we obtain optical gaps of 2.37 and 1.83 eV, respectively, for the barrier and well layer materials. The conduction and valence band offsets, when these two layers are put together, have not been measured. According to Karcher et al. [18], the shift in the conduction band of a-SiN x : H with increasing x is not significant, when x < 1.2 and the a-SiN x : H is prepared by sputtering or with a glow discharge. Hirose and Miyazaki [19] reported that the conduction band of a-SiN0.2 : H only shifts 0.03 eV relative to that of a - S i : H when both are made with a glow discharge. This shift suggests that the offset in the conduction band in the present superlattices is much less than that in the valence band. 3.2. The superlattice structure Figure 2 shows a TEM photograph of the cross-section of a superlattice, which consisted of a total of 36 layers and fabricated to have the same barrier and well widths, both equal to 6 nm.

Fig. 2. The T E M micrograph shows a cross-section of a superlattice structure designed to contain eighteen 6 nm thick barrier layers and eighteen 6 nm thick well layers.

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Fig. 3. This SEM photograph shows layers exposed by ion beam etching a hole in the same superlattice structure as that shown in fig. 2. The etched sample has been coated with about 200 A of gold to reduce sample charging during the observations. This secondary electron image shows that the well and barrier layers are etched at different rates.

The uniformity of the layers as well as the boundaries between the layers are shown. Low angle X-ray diffraction for the sample designed to have 1.5 nm thick wells and 3.0 nm thick barriers confirms the superlattice structure and indicates the period for this superlattice was 4.9 nm. Figure 3 shows a SEM micrograph of the sloping surface of a hole made by ion etching a portion of the same superlattice as that represented in fig. 2. The appearance of the image in the SEM micrograph is due to the surface roughness of the sloping edge of the sputtered hole. The roughness was introduced by the differential ion etching of the two different layer materials. By comparison, a single layer film etched in the same way shows a smooth featureless surface at the etched edge. Thus, the image in the SEM micrograph represents the superlattice layer structure. The period spacings shown in fig. 3 and the known layer thickness indicate that the angle the etched surface makes with the substrate decreases toward the bottom of the hole and is about 0.8 ° at the free film surface. This angle was much smaller than the edges of the hole cut in the etching mask, which implies that most of the etching was accomplished by ions scattered from

S. Zhang et al.

/ a-SiNo 58 : H / a-SiNo 72 : H superlattices

the sloping edges of the holes in the etching mask. Although the SEM photograph only shows part of the multilayer structure and the image is not as clear as the T E M picture, the sample preparation needed for SEM observations is much simpler than that needed for T E M imaging and the former may be useful, in general, to verify that a layer structure is obtained.

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3.3 Optical gaps in the superlattice Figure 4 shows the dependence of the activation energy, Ea, and half of the optical gap, Eo/2, on the well width with the barrier width fixed at 6 nm. The infinite well width represents the result for a single well layer. The optical gap (o) remains virtually constant for well widths larger than 3 nm. However, as the well width is decreased below 3 nm, E o increases. This variation is similar to that reported for a - S i : H / a S i N ~ : H superlattice structures, but the change occurs at well widths of 3 nm rather than the 5 nm reported for a - S i : H / a - S i N ~ : H superlattices [1,4]. This curve could be described by a Kr6nigPenney model using a suitable charge effective mass, although this is not done here. If these were the only observations, they could be explained by a possible quantum size effect. To decide whether this blue shift in E o is caused by an artifact of the Tauc plot, or whether

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least squares fit of the top 19 datapoints was use to draw the straight line.

a Cody plot is more appropriate for determining the optical gap (see ref. [3]), both Tauc (o) and Cody plots ( * ) are shown in fig. 5 for a superlattice with 1.5 nm thick wells and 3 nm barriers. The Tauc plot exhibits a better fit to a straight line for a wider range of the data. In addition, if one uses a Cody plot to deduce the optical gap for a single a-SiNo.ss : H layer, it displays excessive curvature and would suggest an optical gap near 1.2 eV. This number is unacceptably small. These results suggest that a Cody plot is not appropriate for our samples. A series of superlattices were fabricated with 1.5 nm well widths and barrier widths varying from 1.5 to 12 nm. The dependence of the optical gap and the electrical activation energy on barrier width is given in fig. 6, where an infinite barrier width represents an a-SiN0.72 : H film. The optical gap (0) remains unchanged for barrier widths larger than 6 nm but it decreases for barrier widths of = 5 nm and less. The optical gap for a superlattice with a 1.5 nm thick barrier is about 0.3 eV smaller than for one with 6 nm thick barriers.

S. Zhang et aL / a-SiNo.58 : H / a-SiNo.72 ."H superlattices

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3. 4. Electrical properties The room temperature dark conductivity of the well layer material (a-SiN0.58 :H) is about six orders of magnitude larger than that for the barrier layer material (see fig. 7). Hence, the conductivities and the activation energies would be determined by the well layers in these superlattice structures. The dependence of the activation energy on the well width is given in fig. 4 ( * ). The activation energies for the superlattices with 9 and 6 nm well layers are smaller than those for a well layer film alone and they are also smaller than half of the typical gap. As the well width . . . . . . . . . . .

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Fig. 7. The d e p e n d e n c e of the room temperature dark conductivity, o'Rx, and log(trph/O'RT) on the well width. The barrier width was fixed at 6 nm. The lines are drawn as a guide for the eye.

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Fig. 6. The d e p e n d e n c e of E a ( * ) and E o / 2 ( o ) on the barrier layer thickness is shown. The well widths were fixed at 1.5 nm. The line is drawn as a guide for the eye.

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Fig. 8. A plot of log [conductivity] vs. barrier width showing the comparison between the measured values ( * ) and those calculated using eq. (2) ( o ) assuming the fixed prefactor value of 3000 S / c m , the same value as that for a single well layer. The line is drawn as a guide for the eye.

decreases, E a approaches E o / 2 , i.e., the superlattice becomes intrinsic-like. Figure 7 gives the dependence of CrRT and the ratio of %h/~rRT on the well width for the series of samples represented in fig. 4. ~r~ is the photoconductivity of the superlattice measured at room temperature when illuminated with 60 ~ W / c m 2 from the mercury light source. A maximum in ¢rRX is observed at a well width near 6 nm, which can be attributed to charge transferring from the barriers to the wells. The decrease in the dark conductivity for thin well layers is mainly caused by the increase in the activation energies in these layers. O'ph//OrRT increases with decreasing well width for well widths < 2 nm. However, %h varies by about one order of magnitude in this series of samples and it, too, has a maximum value near well widths of 6 nm. The increase in the ratio shown in fig. 7 is due mainly to the large decrease in gRT" Figure 8 shows a plot of log O'RT VS. barrier width where the symbols give the measured values and the o symbols are values calculated using the formula OrRT = O"0 e x p ( - E a / k T ) ,

(2)

where ~r0 = 3000 S/cm; the value for the prefactor obtained from single well layers and the EaS are those given in fig. 6. Satisfactory agreement

S. Zhang et al. / a-SiNo.ss : H / a-SiNo. 72 : H superlattices

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ing photocurrent, Iph, can be described by the equation

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Log [ T i m e ( m i n . ) ] Fig. 9. A typical photocurrent decay curve for the superlattice structures. The dashed line gives the longer term decay rate.

(3)

where t is time and a is the slope of the decay curve and has the value of 0.46 in this case. Other superlattices and the single layer films also give numbers for a close to this value. There is no indication that the degradation in the superlattices is very different from that of the single layers. Following the irradiation of a sample, the photoconductivity can be restored by annealing it at 230°C for about 1 h. Figure 10 shows a plot of log (Conductivity) vs. 1/T for two superlattice structures both before and after exposure (at room temperature) to 12 h of irradiation using the mercury light source. The open circles represent the values obtained before irradiation while the star symbols and the line

6

between the measured and calculated values indicates that o-0 is a constant, suggesting that the same transport mechanism is present in this series of samples. This similarity seems reasonable because the electrical conduction is dominated by the well layers. This result is different from that for the superlattices with varying well widths where a constant value for o-0 was not observed. The activation energy (* symbols in fig. 6) follows Eo/2 rather well as the barrier width decreases and the Fermi level remains near midgap, further suggesting that the same conduction mechanism is operative in this series of samples.

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3.5. Photocurrent degradation and persistent photoconductivity (PPC) All single layers and superlattice samples reported here were photoconducting, and this photoconductivity degraded during illumination with light from the mercury source. Figure 9 shows the photocurrent decay process measured at room temperature, for the superlattice with 1.5 nm thick wells and 3 nm thick barriers. The degrad-

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Fig. 10. The change in the dark conductivity that occurs as a result of irradiating two multilayer structures with the same mercury light source. T h e o symbols were obtained for annealed samples; the solid line and the * symbols show that persistent photoconductivity is present after the multilayer structures are irradiated for 12 h with the mercury light source. The well layers were 1.5 n m thick and the barrier layer thicknesses are as indicated.

228

s. Zhang et al. / a-SiNo.58 : H / a-SiNo 72 : H superlattices

were obtained after irradiation. All data were recorded as the temperature increased. An increase in the conductivity of the superlattice with 12 nm thick barrier layers following irradiation is readily observed, and this has been referred to as persistent photoconductivity or PPC (see, for example, ref. [20]). The magnitude of the PPC is reduced when the barrier width is decreased (e.g., to 1.5 nm) and no PPC is observed in our single layer films. The magnitude of the PPC increased with increasing exposure time up to the maximum exposure time of 12 h. After annealing a multilayer sample at 230°C for 1 h, the dark conductivity was restored to the original value.

4. Discussion When the barrier width decreases, the wavefunctions for quantized states in two adjacent wells could overlap and subbands would form. Miyazaki et al. [4] have used the Kr6nig-Penney model to determine the conditions for the formation of subbands as well as some properties of these bands in a - S i : H / a - S i 3 N 4 : H structures. They concluded that the critical barrier width is 1 nm, i.e., above this value the energy levels in the well are discrete and below this value subbands develop. These subbands spread in energy as the barrier width decreases further and the band width for the first or lowest subband is of the order of 1 meV wide. Although we have not investigated superlattices with barrier widths < 1 nm, we have observed that the optical gap decreases by 0.3 eV as the barrier width decreases from 6 to 1.5 nm (see fig. 6). This reduction in the value of the optical gap is more than two orders of magnitude larger than what would be expected due to the subband bro~tdening suggested by the Kr6nig-Penney model, and implies that the quantum size effect (wavefunction overlap from adjacent wells) cannot explain this result. A similar optical gap decrease with decreasing barrier width has been observed by Bernhard et al. [6] in a - S i : H / a - S i G e x : H multilayers. Their experimental results show that the optical gap is reduced by about 0.1 eV, when the barrier width is decreased from 5 to 1 nm for 1.25 nm thick

well layers. This reduction was explained using coherent multilayer thin-film optics and assuming the presence of transition layers. They attribute this decrease to the absence of a sharp decrease in the absorption coefficient for amorphous semiconductors. In fact, the effect of the absence of a sharp absorption edge on the optical gap differs from that introduced by transition layers. In the former case, the observed increase in E o with increasing barrier width is caused by the measuring technique. When the barrier layer is much thicker than the well layer, the absorbance, a t ( a is the absorption coefficient and t is the thickness), of the total thickness of all of the barrier layers, measured by the spectrophotometer, cannot be neglected in the region used to determine the optical gap E o. This effect leads to a larger observed Eo, while the actual E o for the wells does not change as the well or barrier layer thicknesses are varied. In the latter case, i.e., when a transition layer is present, the optical gap in the wells can change with different well or barrier widths. In our case, it is believed that such a transition layer is present which might be introduced during the deposition process (a) by s o m e mixing of extra N into the well layer and out of the barrier layer as a result of the ammonia ion bombardment, and (b) as the films are deposited, some diffusion of N that is physisorbed or loosely bound may occur at the layer interfaces and this could be enhanced by the ion bombardment. The small differences that can be observed in the thickness between the well and barrier layers in the T E M micrograph of fig. 2 is consistent with such a transition layer hypothesis. For a thick barrier layer and a thin well layer, any extra nitrogen entering the well layer would increase the optical gap of that material. However, the amount of nitrogen that can be introduced into the well is reduced when the barrier 'layer is thin. In other words, both the increasing and the decreasing optical gap with decreasing well and barrier layer thickness, respectively, could be explained by a varying nitrogen content in the well layers. The presence of transition layers is also suggested by the fact that the curve for the optical gap in fig. 6 could be extended to pass near the values ob-

S. Zhang et aL / a-SiNo.s8 : H / a-SiNo.72 : H superlattices

tained for the single well layers and by the fact that lines for O-RTand Orph/OrRXcould pass through the value for the single barrier layer (see fig. 7). Directly verifying the presence of a transition layer by observing the variation in N N across the interface regions is difficult and this has not been done. We have attempted to observe N diffusion into the well layers in a superlattice structure with 2 nm thick wells and thick barrier layers. The results show that no change in E o occurs in a completed superlattice structure for annealing temperatures up to 250°C. The optical gap does not change with well thickness (fig. 4) when the well width is > 3 nm. In these thick well layer superlattices, the carriers (assumed to be electrons) collect in the well layers and this accumulation of electrons shifts the Fermi level above the intrinsic level (see the values for E a in fig. 4). When the well width is reduced to small values ( < 3 nm as shown in fig. 4), interface states may pin the Fermi level near mid-gap. An alternative explanation is that the Fermi level moves to mid-gap due to the increased N concentration in the wells, since aSiN~ : H with a larger N concentration becomes more intrinsic-like. However, since the conductivities for the thinnest-well superlattices are still three orders of magnitude larger than that for a single barrier layer (see fig. 7), the well conductivities still dominate the transport in these samples, and the increasing electrical activation energy for the thinnest well layer is consistent with an increasing band-gap in the well layers. A similar change in the activation energy with well width was also observed in a-Si : H / a - S i N x : H superlattices [2]. In a-Si : H/a-SiN~ : H superlattices, the photoconductive decay which occurs during illumination is normally attributed to the Staebler-Wronski effect and H is believed to play a role [20]. A photoconductivity degradation has been reported in non-hydrogenated a-SiN x by Zhang and Brodie as well as in a-SiN x : H by Kanicki et al. [22]. These latter studies suggest that H may not be involved in the degradation mechanism in aSiN x :H. In addition, the same degradation rates noted in the superlattice structures and the single layer films suggest that neither interface states

229

nor carrier confinement in the wells influence the degradation mechanism very much. The increase in the PPC with increasing barrier width evident in fig. 10 was also observed in the a-Si : H/a-SiNx : H multilayers by Hamed and Fritzsche [20]. They reported that the PPC at room temperature could be five orders of magnitude larger than the original dark conductivity when 80 m W / c m 2 radiation was used, and this PPC is larger than we have observed in any of our samples using 60 p~W/cm 2. It is possible that the maximum PPC was not observed in our samples. Hamed and Fritzsche [20] suggested that some photoexcited carriers in the wells hop through localized states to the barrier layer, where metastable states are produced during irradiation and remain there after the light is turned off. A new Fermi level position is established by these metastable states which occur at different energies in the two layers, and the electron density in the well layers increases. Since no PPC was observed in any of the single well, or barrier layers reported here, the PPC must be related to the the multilayer structure. We assume that a spacial separation of electrons and holes is necessary, but the creation of metastable defects in the two layers must play a major role.

5. Conclusions

Amorphous superlattice structures of a-SiN x : H / a - S i N y : H have been prepared by a new method. The nitrogen content in the silicon nitride layers is varied easily by varying the ammonia ion beam current and leaving all other deposition parameters unchanged. In the a-SiNo.58 : H/a-SiNo.72 : H superlattices, the optical gaps, determined from Tauc plots, vary with both the thickness of the well and the barrier layers when the well or both of these layers are thin. Both of these observations can be explained by a non-abrupt well-barrier interface where the N concentration gradually changes from the well layer to the barrier layer and it is assumed that this changes the NN in the well layers.

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Like a-Si : H / a - S i N x : H, the a-SiN0.s8 : H / a SiN0.72:H superlattices, as well as the single aSiN x : H layers, display a reversible p h o t o c o n d u c tivity d e g r a d a t i o n w h e n i r r a d i a t e d with a m e r c u r y light source. T h e n o n - d e g r a d e d state for these samples can be restored by a n n e a l i n g the samples at 230°C for 1 h. I n o u r samples, only the superlattice structures, a n d not the single layer films, displayed a detectable PPC. This would suggest that a spacial s e p a r a t i o n of electrons a n d holes is r e q u i r e d for P P C in o u r films. T h e a u t h o r s acknowledge with thanks the help of M r F. P e a r s o n at M c M a s t e r University in p r e p a r i n g the T E M samples a n d Drs M a r s h a Singh a n d Bob S h a n n o n at Q u e e n s University for doing the low angle X-ray diffraction e x p e r i m e n t s o n the samples. This work was sponsored by the N a t u r a l Sciences a n d E n g i n e e r i n g R e s e a r c h Council of C a n a d a .

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