Structural, optical and photoelectronic properties of improved PECVD a-Ge:H

]OURNA L OF

Journal of Non-Crystalline Solids 137&138 (1991) 803-808 North-Holland

l g-CRNgI E

Section 14. Amolphous germanium (a-Ge:H) and SiGe alloys STRUCTURAL, OPTICAL AND PHOTOELECTRONIC PROPERTIES OF IMPROVED PECVD a-Ge:H William PAUL Division of Applied Sciences, Harvard University, Cambridge, Massachusetts 02138, USA

The preparation conditions by PECVD for a-Ge:H with photoelectronic properties approaching those of a-Si:H are described. Emphasis is placed on the identification by TEM, SEM, NMR and other techniques of structural defects on a scale of 1 nm and larger, and their elimination by adjustment of deposition parameters. Evidence concerning the valence and conduction band densities of states is reviewed, and preliminary information on the extension of this approach to improvement of a-Si~_=Ge~ :H alloys examined. 1. INTRODUCTION It is generally agreed that the photoelectronic properties of a-Sil_=Gex:H are poorer than those of a-Si:H? Given that (1) the alloys are usually prepared under conditions which optimize a-Si:H, that (2) they risk the complication of spatial fluctuations of the GelSi ratio, and that (3) the fundame0tal processes of transport and phototransport are not well understood even in a-Si:H, various groups have undertaken the study of the pure a-Ge:H end-point. Material with good photoelectronic properties has now been reported. Despite some differences in detail, the general conclusions concerning preparation parameters and optimized photoelectronic properties appear to agree. Our own results 2 are quite close to those of the Siemens group? Several of our papers ~ have described the conclusion from many structural studies that there is a twophase network of high atomic density islands and lower density connective tissue in this class of materials. Others have focussed on the photoelectronic properties and have attempted to relate them consistently to the structure. In a recent review, ~ we suggested that the optical properties were attributable to islands of different structural quality, but that the transport and phototransport were heavily influenced by the amount of highly defective tissue connecting islands. Here we attempt a general assessment of (1) our understanding of the medium range structural order on a scale of 1 nm and greater, (2) the modifications of the PECVD preparation parameters found efficacious in producing a 0022-3093/91/$03.50 © 1991 - Elsevier Science Publishers B.V.

material of low dangling bond and low void content, (3) the evidence for a band structure with reduced density of states in the band tail and energy gap, (4) the relation of our band parameters to those of other investigators and also to extrapolations from the immense catalog of work on a-Sil_=Ge.~:H alloys, and (5) desired improvements on which future work might usefully concentrate. 2. STRUCTURE We seek to produce homogeneous a-Ge:H with a negligible amount of low density connective tissue. (1) On a scale of about 2 nm and larger, the 2-phase island/tissue structure is often revealed by TEM Films of a-Si:H and a-Ge:H with no TEM microstructure on this scale have both been made by us, but the deposition conditions are quite different? ,2 Figure 1 illustrates (a) a-Si:H deposited at 250°C from pure Sill4 at a low deposition rate under conditions minimizing substrate bombardment by ions, (b) a-Ge:H deposited from (GeH4+H2) under similar conditions to (a). Heterogeneity is clearly visible, and (c) a-Ge:H deposited from (GeH4+H2) at an electrode at a plasma-caused self-bias of -150 V with respect to ground, which is expected to promote ion bombardment. No structure is visible. SEM pictures of cross-sections of thick films deposited underthe same conditions as for films (a)-(c) verify the likelihood of columnar structure when heterogeneity is visible in TEM. It is necessary here to discuss briefly the geometry of ourtwo reactors? They have two electrodes differing All rights reserved.

804

W. Paul / Structural, optical and photoelectronic properties of improved PECVD a -Ge:H

in area by x2. An r.f. power of 30 W produces a bias

to correspond to the evolution of contaminant gases

potential at the smaller electrode (the powered cathode, C) of -158 V and at the larger (the anode, A) of

which enter the open-network structure after deposition, and a very high temperature one corresponding

-18 V. All of the films of a-Ge:H and a-Sil_xGex:H with featureless TEM pictures have been deposited at the

to the evolution of GeOx and H2. Differential scanning

smaller electrode of either reactor. We must recognize that we are limited in our description of the precise plasma conditions. Figure 2 shows TEM micrographs of a-Ge:H films deposited at the cathode for spacing of the electrodes increasing from 1.0 to 3.2 cm 4. Although the biases on both electrodes change by only a few percent, the visible film heterostructure increases markedly, and the photoelectronic properties deteriorate. Although the plasma conditions near the cathode must have changed, the changes in the film are clearly

electron diffraction as-deposited, crystallizes close to

calorimetry shows that the C material, amorphous by 430°C in a discontinuous transition s reminiscent of explosive crystallization. The A material crystallizes in 1 or 2 stages depending on the temperature of deposition and the substrate used. 5 Sequential crystallization at temperatures which may differ by 100°C, with a total heat of crystallization closely matching that found in single event crystallization, is a clear (and new) signal of a 2-phase network. The microcrystals formed after the first incident of crystallization are of the order of

not caused by a change in ion bombardment energy. This result differs from the conclusions of the Siemens

8 nm diameter; their properties remain to be explored?

group for their reactor.

than about 2 nm.

(2) We now consider structure on a scale smaller Very detailed information on the

The structure observable in TEM is verified in sev-

environment of H atoms may be obtained from anal-

eral kinds of less direct structural studies. An example

ysis of deuteron magnetic resonance spectra (DMR),

is that of typical infrared absorption spectra for cathode

studied by Professor Richard Norberg on our a-Si:D,H,

(C) and anode (A) films. In addition to the monohy-

a-Sil_xGe~:D,H, and C and A a-Ge:D,H produced by

dride stretch mode at 1875 cm -1 and the wag mode at 560 cm -1 seen in environmentally-stable C films, the

substituting D2 for H2 in the preparation plasma. ~ At least six environments for D (or H) are identifiable:

A films have a stretch mode at 1975 cm -1, bending

(1) tightly-bound Ge-D, (2) D2 or HD molecules in high

modes at 760 and 830 cm -1, and show contamina-

density, as in solid D~. These may occupy the tis-

tion by C and O which increases with exposure to the

sue regions visible in TEM, but also regions smaller

atmosphere. Evolution of 14 on heating a sample at

in size.

a controlled rate may be used to estimate the total

in very small voids or even interstitial positions in the

H-content and also to provide information on how the

fully-coordinated CRN. (4) germanyl rotors, GeDxH3_x. (5) weakly-bound D, giving a DMR signal identifiably

H is bound.

In a-Si:H there are typically two peaks:

(3) isolated D2 or HD molecules, probably

ness, indicates that there are H2 molecules present

different from molecular species and from tig htly-bound Ge-D and (6) para-D2 molecules. Comparison of the

which can leave the film easily, while a peak whose

line-shapes of the DMR spectra for C and A a-Ge:D,H

a low temperature one, independent of sample thick-

(higher) temperature position depends on thickness is

verify that there is more disorder in the environment of

a signature of atomic H from Si-H bonds which exit

Ge-D bonds than Si-D, and more disorder in the Ge-D

by diffusion, with occasional trapping at Si dangling

bonds for the A than for the C Ge. Several new and

bonds. The results for a-Ge:H are more complex. Our

significant conclusions emerge. One is that D2 and HD

C material also shows two major peaks, but the H2

exist in the C material both in the dense and isolated

evolving at low temperature can exist in different sites only revealed by deuteron magnetic resonance (DMR) (see below). The A material may have two additional

forms. Thus this material must have void regions, invisible in TEM, which are of larger dimension than those in our state-of-the-art a-Si:D,H. This is consistent with

peaks: a very low temperature one verified by an RGA

the observation of much low temperature gas evolution

805

W. Paul / SO~ctural, optical and photoelectronic properties of improved PECVD a-Ge:H

(a)

(b)

(c)

FIGURE 1 TEM micrographs for (a) a-Si:H (b) A a-Ge:H and (c) C a-Ge:H. See text for details.

(a)

(b)

(c)

FIGURE 2 TEM micrographs for C a-Ge:H deposited at several electrode spacings D (a) D = 1.0 cm (b) D = 2.0 cm (c) D = 3.2 cm.

in our C films and the conclusion from a comparison of gas evolution, forward recoil scattering and infrared absorption measurements that the total H-concentration is about twice the bonded component. Another important conclusion is that of the existence of weakly-bound single D at the 10 -4 concentration level. While to date this has been investigated mostly in a-Si:D,H films, its presence in a-Ge:D,H also is considered highly likely. We conclude from the DMR studies that even the C aGe:H has void structure on a scale below that seen in TEM or appreciated in GE, IR and DSC investigations. This void structure has also been verified by SAXS, which indicates that the C material has many fewer voids than the A, but that the void distribution in it is much more extensive than in state-of-the-art a-Si:H. 7 Thus it is evident that structural improvements on the scale below 2 nm are still necessary. 3. COMMENTS ON PECVD PARAMETERS These results establish that, in our reactor, different preparation conditions are required for "good" aGe:H than for a-Si:H. This agrees with the conclusions of other groups which use both PECVD and sputtering. However, it appears that there may be different combinations of parameters which all produce reasonable material. Matsuda and Tanaka very early recommended heavy H2 dilution of the plasma. 8 We have found that attempts to produce a-Ge:H from undiluted GeH4 result in dust rather than a coherent film, and that dilution with H2 gives films of good quality. However, we have also produced good quality C material from (GeH4+He) plasmas, which suggests it may be the di-

lution rather than the specific addition of H2 which is responsible for the improvement. Karg et aL 3 have correlated photoelectronic quality with an ion bombardment parameter. We too have found that increasing r.f. power leads to increased substrate bias at the cathode and also to improved photoelectronic properties over the range of power used by Siemens. However, the improvement in properties was found to pass through an extremum and to decrease as the power increase was continued. Moreover, we have found improvements (see Figure 2) in the structure and photoelectronic quality of our films when the substrate self-bias remained unchanged. We conclude that the geometry of the apparatus and the distribution of fields in it are important, as well as the gas and power parameters. Until studies of the plasma and the substrate-piasma reactions are completed, it appears that empirical optimization of individual apparatuses will be needed. Nevertheless, the departures from Sill4 conditions are a guide: while for a-Si:H it is preferred to keep ion bombardment to a minimum, for a-Ge:H it appears to be better to arrange the geometry and increase the power so as to promote bombardment, and to add H2 or He. (The function here of H.2, or HF, may be to etch weakly-adhering material, but this is not proven). Recently, Doyle etal. have confirmed that there are major differences in the radical interactions in GeH4 and Sill4 discharges, and they have proposed tentative reasons why these differences lead to poorer quality a-Ge:H films. 9

W. Paul~Structural, optical and photoelectronic properties of improved PECVD a-Ge:H

806

TABLE 1 Summary of properties of a-Ge:H films deposited at the cathode and anode at 150°C. The illumination used to determine zXI was 8 x 10is photons/cm2sec at 1.25 eV. Sample

Eo4 (eV)

E03 (eV)

r~(2.0 tim)

r/fiT- (cm2/V)

zXI/Id

E,, (eV)

~o (~ cm)

Cathode

1.25

1.10

4.03

3.0x10 -7

1.1 x l 0 -1

0.62

1.0x10 s

Anode

1.24

1.07

3.78

4.2x 10 -1°

1 . 0 x l 0 -~

0.52

3.1 xl03

Sample

Eo,Pos (meV)

o~(0.7eV) (cm -1)

EpL (eV)

z~EpL (eV)

Ns (cm -3)

C,v,zR (at. %)

stress (kbar)

Cathode

51

8.3

0.81

0.19

5x1016

6.4

+4.0

Anode

89

91

unmeasurable

unmeasurable

6 x 1017

5.6

-1.5

It is plausible to extrapolate that the best quality

films the PDS-determined values lie between 42 and

a-Sil_,Gex:H alloys cannot be prepared from a sin-

61 meV and the CPM-determined ones range from 40

gle plasma, and that gentle deposition techniques (re-

to 55 with 15 of the 16 below 50 meV. Thus the va-

mote plasma CVD, photo-CVD, hot-wire) may not be

lence band tails are as sharp as in a-Si:H. This fits

as appropriate for Ge as for Si. These inferences may

earlier reports that Eo is roughly constant through the

change when we understand better the roles of plasma

a-Si~_,~Gex:H alloy series; we have recently also con-

interactions and substrate-plasma processes. Control

firmed this conclusion for our a-Sij_,:Ge,::H alloys at

of the surface mobility of depositing radicals is clearly

the Ge-rich end, deposited under conditions similar to

important.

those used for a-Ge:H. The larger values of £o in the A films are attributed primarily to a deterioration of the

4. PHOTOELECTRONIC PROPERTIES The photoelectronic properties, reported at greater

island material and secondarily to the poorer tissue.

length elsewhere, 2 consistently confirm that mate-

The subband-gap absorption coefficients at 0.7 eV for C films are about 10 cm -1 by PDS and x 2-3 smaller

rial with less heterostructure has sharper absorption

by CPM measurements. This is consistent with the

edges, smaller Urbach tail coefficients, less subband-

poorer a-Ge:H structure below 2 nm and is a factor 10

gap absorption, better photoluminescence and supe-

larger than in good a-Si:H. (5) The transport parame-

rior qffT products. Table 1 lists an earlier comparison

ters E,, and ,7o are not very informative regarding the

of C and A films which we shall use as illustration here,

differences between C and A material, yet there is a

with some addition of statistics: (1) For typical C and

signature, not yet understood, which consistently dis-

A films deposited under identical conditions, the E04's

tinguishes the two: the lno versus

are about the same, implying that the joint valenceconduction band density-of-state (DOS) distributions

material is linear down to 25°C while that for C material is always concave to the 1/T axis, with no sign of

are the same at this energy separation. (2) The differ-

a sharp kink at any temperature. Moreover, the o-0 is

ence (Eo4 - Eo3) is about constant at 0.15 eV for all

always larger in C films for the same E~; we interpret

C films; this is, however, smaller than for the A films,

this to mean that the shift with temperature of the Fermi

and it is also slightly smaller than we found in our ear-

level in C films is much greater, probably because it is sited in a low and rapidly varying part of the gap

lier study of a-Sij_~.Gex:H alloys. See Figure 3. The

;/T relation for A

fact that the upper part of the absorption edge is only a

DOS versus energy relation. (6) The A films have no

little different in C and A materials with distinctly different heterostructures implies that the islands dominate

photoluminescence we could measure, but the C material (38 films) have E(peak) = 0.80±0.02 eV; FWHM ~ E = 0.19±0.02 eV and intensity 10 2 to 10 -3 of

this part of the absorption edge. (3) The Urbach parameter E0 is much smaller in the C material; for 16

a-Si:H. The PL is almost independent of variation of

W. Paul~Structural, optical and photoelectronic properties of improved PECVD a-Ge:H

807

TABLE 2 Correlation of q/.~T with subband-gap absorption coefficient for A a-Si:H, C and A a-Ge:H and a-SiGe:H. Eo4(eV) E(eV) e(cm -~ ) q#T(cm~/V) c~q#T(cm/V) a-Si:H

1.9

1.2

0.5

5x10 -6

2.5×10 -6

C a-Ge:H

1.2

0.7

14(16 films)

2.5x 10-~(16 films)

3.5x10 -6

A a-Ge:H

1.2

0.7

81(6 films)

9× 10-1°(6 films)

7.3x10 -s

a-SiGe:H(1988)

1.4

0.8

20

5x10 -1°

l x 1 0 -s

deposition parameters, and is different from extrapolations from our and Carius 1° earlier measurements on a-Sil_=Gex:H alloys. See Figure 3. (7) The spin density in C material is x 10 smaller than in A material, but x]0 larger than in a-Si:H. This is consistent with the differences in subband-gap absorption. (8) Finally, the most dramatic element in Table 1 is the superior 71/~- of C films. A second illustration involves the photoresponse of C films from reactor 2 at the several electrode spacings used for the TEM's of Figure 2. Figure 4 shows that zM/I (dark) product decreases monotonically as the electrode spacing is increased, which we link directly to the increased heterostructure in Figure 2 (the r e a s o n s for these changes are not relevant here). In both of our reactors the photoelectronic properties change as the deposition parameters are varied. Although this delineates the "best" volume of parameter space for our reactors, it is not directly transferable to others. What is more significant are the correlations of the properties among themselves. Thus, an improved microstructure evidenced by TEM (C films) correlates without exception with decreases in Eo and ~ (0.7 eV) and increases in q#T. Three DMR parameters - - the fraction of tightly-bound D (i.e., H), the fraction of tightlybound D2 or HD, and the inverse relaxation rate, which is proportional to the density of paramagnetic defects, also correlate extremely well with the 7/#T found for identically-deposited films. 6 It is interesting to attempt a correlation between

q#T

and the subband-gap ct., which reasonably represents the subband-gap DOS determining T. The results are shown in Table 2. From the equivalence of a.q#T for aSi:H and C a-Ge:H we infer that the reduction in ~//~7-in

2.0

,

. . . . . .,. . . .

-

• •

1.5

O

E0~ A Z~

~

> •

v

[]

>, 1.0

m

E LU

rq%

Epl"

0.5

AEpl •

00.

•

O@o

¢)d9o

, , , I , , , l l l l l l l l l l l r

0.0

0.2

0.4

0.6

0.8

.0

X FIGURE 3 Dependence on Ge content of Eo.~, E,~3, £p~ and ,AEp~. Solid symbols, 1988; open symbols 1991 0.05

~,ll,~,,i,~,~l~,,,

I

0.04 ,~ 0.03 Q

0.02

•

|

0.01 0 r 1.0

, , , l , , , r l

1.5

....

2.0

I ....

2.5

I

3.0

D (cm) FIGURE 4 AI/I (dark) versus electrode spacing D.

•t

3.5

808

W. Paul~Structural, optical and photoelectronic properties of improved PECVD a-Ge:H

the latter is directly related to the increase in gap DOS represented by c~ (0.7 eV). The much smaller c~71ffT for A a-Ge:H we link to the occurrence of more microstructure (tissue). In the fourth row we have listed data on a-Sil_=Ge=:H from our earlier studies; here too, the reduction in qff-r is greater than expected from (0.8 eV); and again we note that TEM clearly shows the existence of a heterostructure. It is also interesting to attempt a correlation between our present results and extrapolations to pure a-Ge:H of data on SiGe al-

6. CONCLUSIONS (1) Different PECVD conditions are required to prepare a-Si:H and a-Ge:H of comparable photoelectronic quality. (2) Ion bombardment and/or etching appear to be necessary to eliminate heterostructure in a-Ge:H. (3) Structural (and possibly chemical) inhomogeneity on a scale of 1 nm still exists, whose elimination

loys. Aljishi et aL ~ have reported the variation of their

should reduce the gap DOS. (4) Our best a-Ge:H has an absorption edge as sharp as that of a-Si:H, but a gap DOS still x 10 larger,

subband-gap DOS with optical gap for a series of aSil_xGex:H,F alloys. Extrapolation to an optical gap of 1.1 eV (the value for our cathode Ge) suggests a gap DOS of ~ 101~ cm -3 eV -~. This is more than an

which is adequate to explain the lower qy~-. (5) Future work should seek to understand better the DOS in the CB tail, and to extrapolate the methods used here to Ge-rich a-Sil_,,Ge~.:H.

order of magnitude larger than our estimates for our material. Aljishi et al. suggest appropriate parameters for both the entropy model and the kinetic model for defect equilibration which fit their alloy data. Obviously these suggested parameters do not fit our results for pure a-Ge:H. 5. BAND STRUCTURE OF a-Ge:H The similarity of the valence band and valence band tail structures (Eo4, £03, Eo) in C a-Ge:H and a-Si:H is very clear. By contrast, the conduction band tail structure seems to be more complicated, since our time-offlight data taken at 25°C invariably show anomalous log current-log time characteristics. Our EPR experiments suggest an increase in the density of dangling bonds by x l 0 over a-Si:H and, from our crude analysis of ~q#~-, this increase in gap DOS is sufficient to explain the decrease in qff.r. Our structural studies suggest that point and extended defects on a 1 nm scale could be responsible for the increased gap DOS. However, we note also the suggestion by Shu Jin and Ley v-' of an H-related, non-DB defect, possibly a Ge-H-Ge 3-center bond complex. Although an intriguing speculation, one is reminded that the energies of the antibonding states of much simpler Ge-H bonds are assumed with little justification to lie in the conduction band.

ACKNOWLEDGMENTS This work was financially supported by the SERI under Contract XX-8-18131-1. I thank Professor J.H. Chen, Dr. W.A. Turner, Dr. F.C. Marques, Mr. B. Bateman, Mr. S.J. Jones, Ms. D. Pang, Ms. A.E. Wetzel and Mr. P. Wickboldt for their whole-hearted collaboration in our research. REFERENCES 1. An extensive set of references is given in W. Paul et al., Proc. Mat. Res. Soc. 219 (1991) 211, to be published. 2. Reference 1. Also W.A. Turner et ai., J. Appl. Phys. 67 (1990) 7430. 3. RH. Karg et al., J. Non-Cryst. Solids 114 (1989) 477; also, to be published in Solar Energy Materials (1991). 4. P. Wickboldt et al., Phil. Mag., to be published. 5. W. Paul et al., Phil. Mag. B63 (1991)247. 6. Reference 2. Also R.E. Norberg et al., this volume. 7. R. Crandall, private communication. 8. A. Matsuda and K. Tanaka, J. Non-Cryst. Solids 97/98 (1987) 1367. 9. J.R. Doyle et al., J. Appl. Phys., 69 (1991) 4169. 10. R. Carius in Amorphous Silicon and Related Materials, ed. H. Fritzsche (World Scientific, Singapore 1989) p. 939. 11. S. Aljishi et al., in Amorphous Silicon and Related Materials, ed. H. Fritzsche (World Scientific, Singapore 1989) p. 887. 12. Shu Jin and L. Ley, Phys. Rev. B44 (1991) 1066.