- Email: [email protected]
Real-time diagnostics of growth of silicon-germanium alloys on hydrogen-terminated and oxidized silicon (111) surfaces by spectroscopic ellipsometry
Thin Solid Films 343-333 (1999) 477-132
Real-time diagnostics of growth of silicon-germanium alloys on hydrogen-terminated and oxidized silicon (111) surfaces by spectroscopic ellipsometry
Abstract Chemical vapordeposition of amorphous hydrogenated silicon-germanium alloys (a-SiGe:H) was induced by photolysis of disilane and digermane mixtureswith anArF laserat 193nm. The gro\& behavioron hydrogen-terminated andoxide-coveredSiiI 1I) surfaceswas
studiedin realtimein themonoIayerregionby spectroscopic ellipsomerry(1.24.7 eVj. One-,tlvo- andthree-layermodelswereemployedro simuiate the ellipsometric data. For the fiat hydrogen-terminated siricon surface the slow formation of a well-defined monolayer is extracted from the analysis.On the native oxide-coveredsurfacefast initial 3D growth~3s observedwith coalescence at a largerfilm thichess. A rhree-layermode1wasneededto describethe _erowth of ultrathin films after coalescence, wherea third layer, with a thicknessin the nanometer range, takes into account the Lower quality of the ultrathin fiLm layer near the interface. 0 1999 EIsevier Science S.A. All rights reserved. Keywords: Chemical vapnr deposifion; Etiipsometry; Optical properties; Semiconductors
1. Introduction Amorphous hydrogenated silicon-germanium alloys (aSiGe:Hj have attracted much interest for wious optoelectronic applications due to the possibility of adjusting their optical bandgap. Their practical use, however, is so far limited by poor photoelectric properties. The degradation of the fXm quality with ailoying is explained by the existence of inhomogeneities in both composition and SCIWwe. Especially the quality of the interface becomesmofe and more important for applications using ultra-thin films. To come to a better understandingof these problems it is necessaryto investigate the nature and formation of defects and the quality of the interface layer on a molecular level. Plasmadeposition is currently the technique used most commonly to produce this material [l-3]. In these studies the basic structural, optical, and electrical properties of the alloys deposited by plasma-enhanced chemical vapor deposition (CVD) were characterized. Photochemical yapor deposition was also investigated in detail using, for example, an ArF Iaser 1193 nm) to photolyze GeH$$~ mixtures [4] or Ge2HJSi2Homixtures [5]. The growth of
* Corresponding author. Tel,: Jr 49-6221-545205; f&x: f 49-6221544255. E-maii nddress: [email protected] (P. Hess)
solid films by CVD processesis connected with a large number of gas-phase and surface reactions, which are presently not understoodvety well. Considerableprogress has been made recently and can be expected in the near future in elucidating these very complicated scenariosby linear and nonlinear optical probing techniquesasdiscussed in a recent review [6]. Optical probing hasthe advantagethat it can be applied in situ and in real time. It doesnot need vacuum conditions, is usually nondestructke, andprovides specific information on the growing material such as the chemical bonding configuration using FTR spectroscopyin the infrared [7] or optical properties such as the refractive index or dielectric constantsin the yisible and ultraviolet using spectroscopic ellipsometry [S]. Nilthough optical diagnostics is macroscopic in nature. its resuits are Linked to the microscopic properties of the solid surface. In combination with suitable models, optical techniques provide a detailed molecular picmre of the surface and growth processes. Here we report on the real-time diagnosticsof a-SiGe:H growth by spectroscopicellipsometry in the very first stage of surfacereactions. The nucleation processis studiedfor a hydrogen-terminated silicon (I 11) substrateand one with its natural oxide layer. Information on the structure of the first layers is extracted from the ellipsometric measurementsby fitting the data with a sui&bIe model. This reveals, for
OOQI-6090!99/$ - see front matter 0 1999 Elsevier Science S.A. .klI rights reserved PII: SOO10-6090(9SIOl685-X
a-Ge:H
a - SiGe : H
a - Si : H 1
4.0 t
\
1.0 FIN. I. Experimental
set-up
for laser
CVD
and spectroscopic
0.2
ellipsxnerry. Gas
example. the strong dependence of the growth the nature of the substrate surface.
Fi_p. 3. Grouth rafe of a-Gr:H. a-SiCkH. composirion of the deposition ga\ mixture.
Fig. 1 shows a scheme of the experimental set-up with the ArF laser for initiating the CVD process, the UHV chamber with the substrate. and the spectroscopic ellipsometer for real-time diagnostics of the deposition process. The ArF laser induced the CVD process in parallel laser beam-substrate configuration. The laser power was varied between 0.5 and 3 W. At 0.5 W the hydrogen content was about 6%. the SiHJSiH ratio 0.56 and the GeH/SiH ratio 0.1-l for a substrate temperature of 290°C. With increasing laser power and decreasing substrate temperature, the hydrogen content and amount of SiHl groups increased and degraded the electronic properties of the material due to strong distortions of the solid network. All results presented here were obtained at 290°C. The total pressure of the deposition gases and the helium buffer gas was typically 1.5 mbar. The partial pressure of the 7
I
Ge -Ge Ge -Si
800
2
600
kj =
400
x x
200
0
I 200
300 Raman
Fio 3.
and
Raman -.2 the heated
\pecirum hydrogen-free
1
400
for rhe as-grown polycrysralline
shift
500
600
(cm ‘I) amorphous material.
mixturey = @(S&H,)
,
0.6
0.8
/ ((I)(Si,H,)+
1.0
(P(Ge,H,))
process on
2. Experimental
1200
0.4
I
a-Si,,Ge,,.,:H
alloy
and a-Si:H
as a funcrion
of the
precursor specieswas in the range of several microbars. As stated above. the detailed mechanismof the laser-induced gas phase and surface reactions is presently not known. Therefore the silicon/germanium ratio of the film. as comparedwith the gasphasecomposition. was determined by Raman spectroscopy. The deposited amorphous films were heated to 700-800°C lo obtain polycrystalline films without hydrogen- For these films the positions of the SiSi. SiGe. and GeGe peaks.which depend on the composition of the alloy. are available in the literature [9] and could be used for the composition analysis.Fig. 2 gives an example of the spectrumof a specific alloy. The halfwidth of the Raman peaks dependson the structural order of the film material. The comparatively narrow peaks observed for the amorphous material indicate a relatively high shortrange order of the depositedalloys. As shownin Fig. 3, the growth rate in&easedfrom I .75 to 3.8 nm/min on going from a-Si:H to a-Ge:H. The non-linear growth behavior of a-SiGe:H in between can be explained by a higher optical absorption coefficient of Ge:H, [lo] and the strongcrosschemistry betweensilicon- andgermaniumbearing species[J]. There was no indication of a catalytic effect of germanium at 290°C. in agreementwith previous findings [-$I. It has been postulated that the desorption of hydrogen by germanium from GeH-surface bondsacceleratesthe growth processdramatically [IO]. This effect could be due to thermally induced hydrogen desorption from GeH-surface groups at temperatureshigher than 190°C. It should be noted that in the SiGe alloys hydrogen is mainly bonded to silicon. in agreement with previous results obtained by HREELS [I 11. The ellipsometer (Sopra ESJG O.MA) was a rotating polarizer set-up with a high pressurexenon arc lamp as light source and a usable wavelength range between I.2 and 1.7 eV. The light beam was collimated to 5 mm diameter and fed into the vacuum chamber through an entrance window of fused silica (Suprasil I). The angle of
P. Hess.
1. Opahle
/
Thin
Solid
6
3.0 Pf,G!On
3.5 energy [evj
0
4.0
4.5
Fig. ii. Real part TZand imaginary part k of tk complex index of refraction for oxide-free c-Si (dashed line) and silicon wit11 a native oxide layer (solid line).
incidence was between73 and 74”. After reflection from the substratesurface the light travels through Ihe e,xit window and analyzer and is coupled into an optical fiber. which connects the analyzer arm with a prism spectrograph.The spectrum is then projected onto an optical multichannel analyzer (OMA) with 1024 pixels covering the wavelength range given above. The OMA was calibrated against the spectrum of a mercury lamp. For the acquisition of spectra two pixels are grouped together LOcompensate for the charge drifts to vicinal diodes. With 512 points per spectrum, less than two secondswere required to processone spectrum.The OLMA integratesthe incident it-radianceover an exposure time set to a quarter period of the rotating polarizer.
Films
343-344
(19991
427432
‘$29
To exwact information from ellipsometric measurements a model with suitable parametersmust be fitted to the dara, since the film properties cannot be extracted directly from the measuredeilipsograms.In the following a linear regression analysis with Ca-minimization is appIied to determine the parametersof the growth model selectedto describethe main features of the particular deposition process. To interpret the ellipsometric tan(XP(Ei) and cos(il{E)) spectra severai refractive index functions have to be known. For c-Si with native oxide the bulk refractive index can be determined by measuring the substrate at room temperature, and taking the literature values for the T-Si and SiO? dielectric functions into account. A least squaresregressionanalysis and rhe Brtiggemann effective medium approximation (E&M) were applied to obtain the thicknessand void volume fraction of the native oxide overlayer and the actual angle of incidence, which can vary slightly with different substrates.Fig. 4 iilustrates the effect of a natural oxide layer with a typical thicknessof 2.5 nm by comparing the real part 71and imaginary part X- of the complex index of refraction for a pure and native-otidecovered silicon substrate.The silicon spectrum showstwo characteristic features, the direct transitions around E = 3.3 eV (criticai poinls ED’ and El) and the direct transition at E = 4.3 eV (critical point Ez), The band structure of silicon, illustrating the indirect bandgapEg and the direct UV bandto-band transitions Eo', E, and El, is shown in Fig. 5. The substrateis then measuredat the deposition temperature (29O”Cj to obtain the bulk refractive indicesby application of the following equation for the complex index of refraction (N = rr - iX-) N = :vOsin(pO)[I + ((1 - p)/(l + ~))3tan’(~0)]1’z
(1)
with p = fan{ ?E,expiiilj Next. the refractive indes for the depositedfilm is determined at the deposition temperature from a thick opaque layer, again using Eq. (1). These values can be corrected for the influence of the surface roughness,modeled as an overlayer consistingof material with bulk optical constants and voids as describedbelow. 3. Modeling
k=
(O:O,O) r
Fig. 5. Band structure of siiicon illustrating the direct interband transitions Eo’. El and E!.
If the optical constantsare tiown, spectroscopicellipsometry allows the analysis of film growth from nucleation to the point were the film becomes opaque using suitable models. A one-layer model may be used to describe the earIJ nucleation stageup to coalescence(seeFig. 6). The optical properties are obtained by mixing the optical constants of the bulk film material with voids (holes) in a Brtiggemann EMA to describe the density-deficient film. Coalescence occursfor void volume fractions below 05, at Ihe maximum
1.0,. . . ? e 0.8 L
Surface
One
/ , 1 d,
Layer
T:vo Layers
Fig. 6. Definition models.
Three
of the film pzrameterl:
3.5
of the one-. MO-,
_ 3.0 F c
2.5
0.0
a - Si3,Gec3
0
n H-terminated Si (111) o Oxide covered Si (111)
1 L
'= 0
I
3
:H o
3
0 0
50
100
150
a - Si,,Ge, :H n H - terminated Si (111) l Oxide covered Si (111)
:.
.
nnd three-layer
I
L
-
1
Layea
thickness of the rough surface layer with low packing density [ 121. To describe film growth after coalescence a model with two film growth parameters must be introduced. namely the thickness of the rough surface layer d,. with the residual surface roughness remaining after the first monolayer forms. and the film thickness ti, itself, as indicated in Fig. 6 [ 12.131. This model is applicable only if the growing film has the same optical properties for the whole thickness. starting at the substrate interface. since homogeneous film properties are assumed in this two-layer model. Quite often the interface region has a lower quality. Even if this defect-rich interface layer is only a few nanometers thick. it will influence the fitted film properties up to a film thickness of about ‘O-30 nm. In this situation the standard deviation of the fit procedure can be decreased drastically by introducing a third layer with thickness cl: at the interface with an increased void fraction (see Fig. 6). Note that the three-layer model also depends on the assumption of homogeneous layers, Lvhich may be not very realistic for the interface layer. The hydrogen concentration in the a-SiGe:H alloys is also described on the basis of a Briiggemann EMA. by mixing the film properties with voids as in the simulation of the surface roughness. This seems to be a reasonable approximation for hydrogen concentrations below 10%.
4.0
s ‘S
I
200
250
Time [sj Fif. 7. Film thickness of rhr a-Si,:Ge,,,:H film 3s i! function hydrogen-wminated und oxide-covered SIC 11 1) surface.
of rime for a
0.0 0.0
0.5 I
1.0 I Film thickness
1.5 I
2.0 I
2.5 !
[nm]
Fig. 8. Compnrijon of the rwlb of the one-lnyer model for the boid volume fraction of the a-Si;rCe,.l:H alloy as 3. function of film thickness for a hydrogen-wrminared and wide-covered Si( I I I) surfxc.
4. Results As an example. the formation of the a-Si;iGe,l:H alloy is discussed in more detail in the following section. The growth process was monitored Lvith high resolution by recording up to 100 spectra during 1.5 nm growth on a Hterminated Si( 1 II) surface. This allowed the real-time observation of monolayer formation with submonolayer resolution. On the oxide-covered substrate very fast growth was detected initially, faster than the integral growth rate of the bulk film. as presented in Fig. 7. One reason for this behavior is certainly the much larger roughness of this surface. This high growth rate limited the number of spectra that could be recorded during this stage. On the H-terminated silicon surface the film thickness stopped increasing after about lo-20 s. The resulting plateau lasted for about 10 s. and then a slow increase of the thickness was observed with a growth rate still below the integral rate. A similar behavior was reported before for the growth rate of a-Ge:H, however. with a much smaller time resolution of optical detection [S]. In this case it was concluded, on the basis of the measured optical constants. that alloying between silicon and germanium occurred in the first layer on the hydrogenated silicon surface but not on the oxide-covered silicon surface. The initial growth behavior ws modeled with the single-layer model. The optical properties were simulated by mixing the properties of the bulk film material with voids using the EMA. It should be noted that clusters may have a different band structure at the nucleation stage, and thus diRerent optical properties. These effects will be discussed in more detail later and have not been taken into account in the present simulation. Nevertheless, the one-layer model yields the following interesting features: For the H-terminated silicon surface the thickness
”
P. Hess, I. Opahle / Thin Sdid Fiirns 343-344
0.0 0
$ ! / , t , < 20 40 60 eo Flln thickness
[m]
0.0 1, / , 1 0
20
( , , j
40 60 80 F:im thickness [nn]
Fig. 9. Comparison of the Z~T values for the twc- and tier-layer model applied to a-Si3,G~3:H growth on a hydrogen-terminated and oxidecovered Sii 1 I 1) surface.
of the surface layer is essentially constant for some time s and the value of about 3 A correspondsto one monolayer (see Fig. 8). The void fraction decreasesduring this time from over 90% to below 50%, indicating coalescenceof the surface coverage at this point. Obviously a highly ordered first layer is formed on the flat H-terminated silicon surface. The oxide-covered surface did not exhibit this discontinuous behavior of the void fraction. For the first measurementa value of 60% was found and 50% was reached near 10 run film thickness, as can be seenin Fig. 8. This is consistent with three-dimensionalcIuster growth, without formation of a well defined monolayer. This behavior seemsto be controlled by the initial roughnessof the surface, the different lattice parameters. and the chemical nature of the surface bonds. Description of film growth with the one- or two-layer model after coalescenceleads to characteristic deviations between experiment and model. As can be seenin Fig. 9, the larger 2g values, describing the quality of the fitting procedure. extend to film thicknessesof 20 nm for both surfaces. These deviations can be decreasedto the usual level of zrr < 0.01 by introducing an additional thin interface layer with a higher void fraction than in the bulk film. The thickness of this third low quality interface layer depends,asexpected, on the nature of the substratesurface and the deposition conditions. The vaiues of [his additional fit parametervaried between 0.7 and 4 nm, where the smaller interface layer thicknesses were observed for the Hterminated surface and the larger ones for oxide-covered silicon.
5. Discussion The chemicai mechanismfor the deposition of a-SiGe:H is extremely complicated and presently not well understood. Therefore, the SiiGe ratio in the film had to be measuredby Ramanspectroscopy.This analysis indicated thar incorporation of germanium into the film was a factor of five more
11959) 427432
i31
effective than for silicon. This effect was at least partly causedby the higher optical absorption coefficient of digermane for the ArF-laser radiation. Since the exact degreeof optical absorption is presently not known, it is not possible to estimate the contribution of cross chemistry to the preferred deposition of germanium. Pulsedlaser depositionallowed weil-resolved diagnostics of the formation of the alloy during the first stageof growth in real time. This madeit possiblefor the first time to monitor the changesin the growth moderesulting from the differences in Ihe surface structure and morphology aswell as the nature of surface bonding. Despite the large information content of spectroscopicellipsomerry a complete description of such a complicated system can, of course, not be achieved by one technique. In the present work modeling has been oriented towards extraction of information on the film structure and morphology. To reach this goal it was necessaryco estimate the optical constantsby very simpIe assumptionsasdescribedabove. Theseapproximations may lead to errors if, for example, well-defined homogeneous Iayers no longer exist, as in the case of self-organized 3D growth or inhomogeneities. The index of refraction is a property of the soIid intimately connectedwith the band structure. Recently the optical properties of germanium quantum dots with alterage diametersin the range 2.5-13 nm were analyzed by investigating the dependenceof the direct interband transitionsE, and E? on the size of the clusters at 77 K [14]. The main resuItsobtained were a weakening of the intensity of the ET peak and a small blueshift of the E? peak with the reduction of the cluster size. Altogether, however, theseeffects seem to be small in comparison with the shifts and broadening observed in an amorphousmaterial. In conclusion, opticai investigations with spectroscopic ellipsometry offer unprecedentedopportunities to address the moIecular details of surface chemistry and morphology of semiconductorsin the interface region with monolayer resolution. On the other hand, films with thicknessesgreater than several monolayers usually are adequateto simulate bulk-terminated surfaces.Therefore bulk optical constants may deliver indeed a reasonable description of the optical properties of uitra-thin layers.
Acknowledgements Financial support of this work by the Bundesministerium fir Bildung und Forschung (BMBF) under contract no. 0329172A9 and by rhe Fonds der Chemischenlndustrie is gratefully acknowiedged.
References [I]
31. Zeman, (19901215.
I. Fe&a,
M.J. Gecrts, J.W. Metselaar.
Xppi. Surf. Sci. 16
132
P. Hess, I. 0,mlde
/ Tlti11 Solid Films 343-344
:2] V.I. Kuznersov. M. Zeman. L.L..I\. Vosreen. B.S. Guar. J.W. Mersehr. J. Appl. Phys. SO (1996) 6196. [3] Y.-P. Chou. S.-C. Lee. J. Appl. Phys. 83 (1998) 11 I I. [-I] H.H. Burke. I.P. Hennan. V. Tavitirn. J.G. Ed-n. Appl. Phys. Lru. 55 l19S9) 153. [5] C. Li. S. John. S. Banerjee. J. Elecrron. Mater. 3-l I 1996) 875. [6] P. Hess, m: M.D. Allendorf. C. Bernard (Eds.). Chemical Vapor Deposirion. Procredmgs97-25. The Electrochemical Society. Pennington. NJ. 1997, pp. 616. [7] I. Knobloch. P. Hess, Appl. Phys. Lett. 69 (1996) 4011. 181 M. Barth. P. Hess. Appl. Phys. Len. 69 ( 1996, 1710. [9] M.A. Renucci. J.B. Rznucci. M. Cardona. in: M. Balkanski (Ed.,.
f 1999) 4_77,/?_7
Proc. Conf. on Light Scatrcring in Solids. Rammarion. Paris, 1971, pp. 31-6. [IO] B. Meyerson. K. L;ram. F. Lcgoucs. Appl. Phys. Lett. 53 (1968) 2.555. [Ill H. Sasaki. M. Drguchi. K. Sate. $1. Aiga. J. Non-Crysl. Solids I I4 I 19S9) 671. [I?] H.V. ?4guyen, Y. Lu, S. Kim, hl. WAgi. R.W. Collins, Phys. Rev. Lett. 74 (1995) 3880. [ I31 Y.M. Li. I. An. H.V. Nguyen, C.R. Wrunski. R.W. CoIlin>, Phys. Rev. Lcrt. 68 (1992) 1814. [l-l] P. Toanini, L.C. Andreani. M. G&o. A. Srcllu. P. Cheyssac. R. Kof&n. A. Migliori, Phy,. Rev. B 53 (1996) 6992.
Real-time diagnostics of growth of silicon-germanium alloys on hydrogen-terminated and oxidized silicon (111) surfaces by spectroscopic ellipsometry
Abstract Chemical vapordeposition of amorphous hydrogenated silicon-germanium alloys (a-SiGe:H) was induced by photolysis of disilane and digermane mixtureswith anArF laserat 193nm. The gro\& behavioron hydrogen-terminated andoxide-coveredSiiI 1I) surfaceswas
studiedin realtimein themonoIayerregionby spectroscopic ellipsomerry(1.24.7 eVj. One-,tlvo- andthree-layermodelswereemployedro simuiate the ellipsometric data. For the fiat hydrogen-terminated siricon surface the slow formation of a well-defined monolayer is extracted from the analysis.On the native oxide-coveredsurfacefast initial 3D growth~3s observedwith coalescence at a largerfilm thichess. A rhree-layermode1wasneededto describethe _erowth of ultrathin films after coalescence, wherea third layer, with a thicknessin the nanometer range, takes into account the Lower quality of the ultrathin fiLm layer near the interface. 0 1999 EIsevier Science S.A. All rights reserved. Keywords: Chemical vapnr deposifion; Etiipsometry; Optical properties; Semiconductors
1. Introduction Amorphous hydrogenated silicon-germanium alloys (aSiGe:Hj have attracted much interest for wious optoelectronic applications due to the possibility of adjusting their optical bandgap. Their practical use, however, is so far limited by poor photoelectric properties. The degradation of the fXm quality with ailoying is explained by the existence of inhomogeneities in both composition and SCIWwe. Especially the quality of the interface becomesmofe and more important for applications using ultra-thin films. To come to a better understandingof these problems it is necessaryto investigate the nature and formation of defects and the quality of the interface layer on a molecular level. Plasmadeposition is currently the technique used most commonly to produce this material [l-3]. In these studies the basic structural, optical, and electrical properties of the alloys deposited by plasma-enhanced chemical vapor deposition (CVD) were characterized. Photochemical yapor deposition was also investigated in detail using, for example, an ArF Iaser 1193 nm) to photolyze GeH$$~ mixtures [4] or Ge2HJSi2Homixtures [5]. The growth of
* Corresponding author. Tel,: Jr 49-6221-545205; f&x: f 49-6221544255. E-maii nddress: [email protected] (P. Hess)
solid films by CVD processesis connected with a large number of gas-phase and surface reactions, which are presently not understoodvety well. Considerableprogress has been made recently and can be expected in the near future in elucidating these very complicated scenariosby linear and nonlinear optical probing techniquesasdiscussed in a recent review [6]. Optical probing hasthe advantagethat it can be applied in situ and in real time. It doesnot need vacuum conditions, is usually nondestructke, andprovides specific information on the growing material such as the chemical bonding configuration using FTR spectroscopyin the infrared [7] or optical properties such as the refractive index or dielectric constantsin the yisible and ultraviolet using spectroscopic ellipsometry [S]. Nilthough optical diagnostics is macroscopic in nature. its resuits are Linked to the microscopic properties of the solid surface. In combination with suitable models, optical techniques provide a detailed molecular picmre of the surface and growth processes. Here we report on the real-time diagnosticsof a-SiGe:H growth by spectroscopicellipsometry in the very first stage of surfacereactions. The nucleation processis studiedfor a hydrogen-terminated silicon (I 11) substrateand one with its natural oxide layer. Information on the structure of the first layers is extracted from the ellipsometric measurementsby fitting the data with a sui&bIe model. This reveals, for
OOQI-6090!99/$ - see front matter 0 1999 Elsevier Science S.A. .klI rights reserved PII: SOO10-6090(9SIOl685-X
a-Ge:H
a - SiGe : H
a - Si : H 1
4.0 t
\
1.0 FIN. I. Experimental
set-up
for laser
CVD
and spectroscopic
0.2
ellipsxnerry. Gas
example. the strong dependence of the growth the nature of the substrate surface.
Fi_p. 3. Grouth rafe of a-Gr:H. a-SiCkH. composirion of the deposition ga\ mixture.
Fig. 1 shows a scheme of the experimental set-up with the ArF laser for initiating the CVD process, the UHV chamber with the substrate. and the spectroscopic ellipsometer for real-time diagnostics of the deposition process. The ArF laser induced the CVD process in parallel laser beam-substrate configuration. The laser power was varied between 0.5 and 3 W. At 0.5 W the hydrogen content was about 6%. the SiHJSiH ratio 0.56 and the GeH/SiH ratio 0.1-l for a substrate temperature of 290°C. With increasing laser power and decreasing substrate temperature, the hydrogen content and amount of SiHl groups increased and degraded the electronic properties of the material due to strong distortions of the solid network. All results presented here were obtained at 290°C. The total pressure of the deposition gases and the helium buffer gas was typically 1.5 mbar. The partial pressure of the 7
I
Ge -Ge Ge -Si
800
2
600
kj =
400
x x
200
0
I 200
300 Raman
Fio 3.
and
Raman -.2 the heated
\pecirum hydrogen-free
1
400
for rhe as-grown polycrysralline
shift
500
600
(cm ‘I) amorphous material.
mixturey = @(S&H,)
,
0.6
0.8
/ ((I)(Si,H,)+
1.0
(P(Ge,H,))
process on
2. Experimental
1200
0.4
I
a-Si,,Ge,,.,:H
alloy
and a-Si:H
as a funcrion
of the
precursor specieswas in the range of several microbars. As stated above. the detailed mechanismof the laser-induced gas phase and surface reactions is presently not known. Therefore the silicon/germanium ratio of the film. as comparedwith the gasphasecomposition. was determined by Raman spectroscopy. The deposited amorphous films were heated to 700-800°C lo obtain polycrystalline films without hydrogen- For these films the positions of the SiSi. SiGe. and GeGe peaks.which depend on the composition of the alloy. are available in the literature [9] and could be used for the composition analysis.Fig. 2 gives an example of the spectrumof a specific alloy. The halfwidth of the Raman peaks dependson the structural order of the film material. The comparatively narrow peaks observed for the amorphous material indicate a relatively high shortrange order of the depositedalloys. As shownin Fig. 3, the growth rate in&easedfrom I .75 to 3.8 nm/min on going from a-Si:H to a-Ge:H. The non-linear growth behavior of a-SiGe:H in between can be explained by a higher optical absorption coefficient of Ge:H, [lo] and the strongcrosschemistry betweensilicon- andgermaniumbearing species[J]. There was no indication of a catalytic effect of germanium at 290°C. in agreementwith previous findings [-$I. It has been postulated that the desorption of hydrogen by germanium from GeH-surface bondsacceleratesthe growth processdramatically [IO]. This effect could be due to thermally induced hydrogen desorption from GeH-surface groups at temperatureshigher than 190°C. It should be noted that in the SiGe alloys hydrogen is mainly bonded to silicon. in agreement with previous results obtained by HREELS [I 11. The ellipsometer (Sopra ESJG O.MA) was a rotating polarizer set-up with a high pressurexenon arc lamp as light source and a usable wavelength range between I.2 and 1.7 eV. The light beam was collimated to 5 mm diameter and fed into the vacuum chamber through an entrance window of fused silica (Suprasil I). The angle of
P. Hess.
1. Opahle
/
Thin
Solid
6
3.0 Pf,G!On
3.5 energy [evj
0
4.0
4.5
Fig. ii. Real part TZand imaginary part k of tk complex index of refraction for oxide-free c-Si (dashed line) and silicon wit11 a native oxide layer (solid line).
incidence was between73 and 74”. After reflection from the substratesurface the light travels through Ihe e,xit window and analyzer and is coupled into an optical fiber. which connects the analyzer arm with a prism spectrograph.The spectrum is then projected onto an optical multichannel analyzer (OMA) with 1024 pixels covering the wavelength range given above. The OMA was calibrated against the spectrum of a mercury lamp. For the acquisition of spectra two pixels are grouped together LOcompensate for the charge drifts to vicinal diodes. With 512 points per spectrum, less than two secondswere required to processone spectrum.The OLMA integratesthe incident it-radianceover an exposure time set to a quarter period of the rotating polarizer.
Films
343-344
(19991
427432
‘$29
To exwact information from ellipsometric measurements a model with suitable parametersmust be fitted to the dara, since the film properties cannot be extracted directly from the measuredeilipsograms.In the following a linear regression analysis with Ca-minimization is appIied to determine the parametersof the growth model selectedto describethe main features of the particular deposition process. To interpret the ellipsometric tan(XP(Ei) and cos(il{E)) spectra severai refractive index functions have to be known. For c-Si with native oxide the bulk refractive index can be determined by measuring the substrate at room temperature, and taking the literature values for the T-Si and SiO? dielectric functions into account. A least squaresregressionanalysis and rhe Brtiggemann effective medium approximation (E&M) were applied to obtain the thicknessand void volume fraction of the native oxide overlayer and the actual angle of incidence, which can vary slightly with different substrates.Fig. 4 iilustrates the effect of a natural oxide layer with a typical thicknessof 2.5 nm by comparing the real part 71and imaginary part X- of the complex index of refraction for a pure and native-otidecovered silicon substrate.The silicon spectrum showstwo characteristic features, the direct transitions around E = 3.3 eV (criticai poinls ED’ and El) and the direct transition at E = 4.3 eV (critical point Ez), The band structure of silicon, illustrating the indirect bandgapEg and the direct UV bandto-band transitions Eo', E, and El, is shown in Fig. 5. The substrateis then measuredat the deposition temperature (29O”Cj to obtain the bulk refractive indicesby application of the following equation for the complex index of refraction (N = rr - iX-) N = :vOsin(pO)[I + ((1 - p)/(l + ~))3tan’(~0)]1’z
(1)
with p = fan{ ?E,expiiilj Next. the refractive indes for the depositedfilm is determined at the deposition temperature from a thick opaque layer, again using Eq. (1). These values can be corrected for the influence of the surface roughness,modeled as an overlayer consistingof material with bulk optical constants and voids as describedbelow. 3. Modeling
k=
(O:O,O) r
Fig. 5. Band structure of siiicon illustrating the direct interband transitions Eo’. El and E!.
If the optical constantsare tiown, spectroscopicellipsometry allows the analysis of film growth from nucleation to the point were the film becomes opaque using suitable models. A one-layer model may be used to describe the earIJ nucleation stageup to coalescence(seeFig. 6). The optical properties are obtained by mixing the optical constants of the bulk film material with voids (holes) in a Brtiggemann EMA to describe the density-deficient film. Coalescence occursfor void volume fractions below 05, at Ihe maximum
1.0,. . . ? e 0.8 L
Surface
One
/ , 1 d,
Layer
T:vo Layers
Fig. 6. Definition models.
Three
of the film pzrameterl:
3.5
of the one-. MO-,
_ 3.0 F c
2.5
0.0
a - Si3,Gec3
0
n H-terminated Si (111) o Oxide covered Si (111)
1 L
'= 0
I
3
:H o
3
0 0
50
100
150
a - Si,,Ge, :H n H - terminated Si (111) l Oxide covered Si (111)
:.
.
nnd three-layer
I
L
-
1
Layea
thickness of the rough surface layer with low packing density [ 121. To describe film growth after coalescence a model with two film growth parameters must be introduced. namely the thickness of the rough surface layer d,. with the residual surface roughness remaining after the first monolayer forms. and the film thickness ti, itself, as indicated in Fig. 6 [ 12.131. This model is applicable only if the growing film has the same optical properties for the whole thickness. starting at the substrate interface. since homogeneous film properties are assumed in this two-layer model. Quite often the interface region has a lower quality. Even if this defect-rich interface layer is only a few nanometers thick. it will influence the fitted film properties up to a film thickness of about ‘O-30 nm. In this situation the standard deviation of the fit procedure can be decreased drastically by introducing a third layer with thickness cl: at the interface with an increased void fraction (see Fig. 6). Note that the three-layer model also depends on the assumption of homogeneous layers, Lvhich may be not very realistic for the interface layer. The hydrogen concentration in the a-SiGe:H alloys is also described on the basis of a Briiggemann EMA. by mixing the film properties with voids as in the simulation of the surface roughness. This seems to be a reasonable approximation for hydrogen concentrations below 10%.
4.0
s ‘S
I
200
250
Time [sj Fif. 7. Film thickness of rhr a-Si,:Ge,,,:H film 3s i! function hydrogen-wminated und oxide-covered SIC 11 1) surface.
of rime for a
0.0 0.0
0.5 I
1.0 I Film thickness
1.5 I
2.0 I
2.5 !
[nm]
Fig. 8. Compnrijon of the rwlb of the one-lnyer model for the boid volume fraction of the a-Si;rCe,.l:H alloy as 3. function of film thickness for a hydrogen-wrminared and wide-covered Si( I I I) surfxc.
4. Results As an example. the formation of the a-Si;iGe,l:H alloy is discussed in more detail in the following section. The growth process was monitored Lvith high resolution by recording up to 100 spectra during 1.5 nm growth on a Hterminated Si( 1 II) surface. This allowed the real-time observation of monolayer formation with submonolayer resolution. On the oxide-covered substrate very fast growth was detected initially, faster than the integral growth rate of the bulk film. as presented in Fig. 7. One reason for this behavior is certainly the much larger roughness of this surface. This high growth rate limited the number of spectra that could be recorded during this stage. On the H-terminated silicon surface the film thickness stopped increasing after about lo-20 s. The resulting plateau lasted for about 10 s. and then a slow increase of the thickness was observed with a growth rate still below the integral rate. A similar behavior was reported before for the growth rate of a-Ge:H, however. with a much smaller time resolution of optical detection [S]. In this case it was concluded, on the basis of the measured optical constants. that alloying between silicon and germanium occurred in the first layer on the hydrogenated silicon surface but not on the oxide-covered silicon surface. The initial growth behavior ws modeled with the single-layer model. The optical properties were simulated by mixing the properties of the bulk film material with voids using the EMA. It should be noted that clusters may have a different band structure at the nucleation stage, and thus diRerent optical properties. These effects will be discussed in more detail later and have not been taken into account in the present simulation. Nevertheless, the one-layer model yields the following interesting features: For the H-terminated silicon surface the thickness
”
P. Hess, I. Opahle / Thin Sdid Fiirns 343-344
0.0 0
$ ! / , t , < 20 40 60 eo Flln thickness
[m]
0.0 1, / , 1 0
20
( , , j
40 60 80 F:im thickness [nn]
Fig. 9. Comparison of the Z~T values for the twc- and tier-layer model applied to a-Si3,G~3:H growth on a hydrogen-terminated and oxidecovered Sii 1 I 1) surface.
of the surface layer is essentially constant for some time s and the value of about 3 A correspondsto one monolayer (see Fig. 8). The void fraction decreasesduring this time from over 90% to below 50%, indicating coalescenceof the surface coverage at this point. Obviously a highly ordered first layer is formed on the flat H-terminated silicon surface. The oxide-covered surface did not exhibit this discontinuous behavior of the void fraction. For the first measurementa value of 60% was found and 50% was reached near 10 run film thickness, as can be seenin Fig. 8. This is consistent with three-dimensionalcIuster growth, without formation of a well defined monolayer. This behavior seemsto be controlled by the initial roughnessof the surface, the different lattice parameters. and the chemical nature of the surface bonds. Description of film growth with the one- or two-layer model after coalescenceleads to characteristic deviations between experiment and model. As can be seenin Fig. 9, the larger 2g values, describing the quality of the fitting procedure. extend to film thicknessesof 20 nm for both surfaces. These deviations can be decreasedto the usual level of zrr < 0.01 by introducing an additional thin interface layer with a higher void fraction than in the bulk film. The thickness of this third low quality interface layer depends,asexpected, on the nature of the substratesurface and the deposition conditions. The vaiues of [his additional fit parametervaried between 0.7 and 4 nm, where the smaller interface layer thicknesses were observed for the Hterminated surface and the larger ones for oxide-covered silicon.
5. Discussion The chemicai mechanismfor the deposition of a-SiGe:H is extremely complicated and presently not well understood. Therefore, the SiiGe ratio in the film had to be measuredby Ramanspectroscopy.This analysis indicated thar incorporation of germanium into the film was a factor of five more
11959) 427432
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effective than for silicon. This effect was at least partly causedby the higher optical absorption coefficient of digermane for the ArF-laser radiation. Since the exact degreeof optical absorption is presently not known, it is not possible to estimate the contribution of cross chemistry to the preferred deposition of germanium. Pulsedlaser depositionallowed weil-resolved diagnostics of the formation of the alloy during the first stageof growth in real time. This madeit possiblefor the first time to monitor the changesin the growth moderesulting from the differences in Ihe surface structure and morphology aswell as the nature of surface bonding. Despite the large information content of spectroscopicellipsomerry a complete description of such a complicated system can, of course, not be achieved by one technique. In the present work modeling has been oriented towards extraction of information on the film structure and morphology. To reach this goal it was necessaryco estimate the optical constantsby very simpIe assumptionsasdescribedabove. Theseapproximations may lead to errors if, for example, well-defined homogeneous Iayers no longer exist, as in the case of self-organized 3D growth or inhomogeneities. The index of refraction is a property of the soIid intimately connectedwith the band structure. Recently the optical properties of germanium quantum dots with alterage diametersin the range 2.5-13 nm were analyzed by investigating the dependenceof the direct interband transitionsE, and E? on the size of the clusters at 77 K [14]. The main resuItsobtained were a weakening of the intensity of the ET peak and a small blueshift of the E? peak with the reduction of the cluster size. Altogether, however, theseeffects seem to be small in comparison with the shifts and broadening observed in an amorphousmaterial. In conclusion, opticai investigations with spectroscopic ellipsometry offer unprecedentedopportunities to address the moIecular details of surface chemistry and morphology of semiconductorsin the interface region with monolayer resolution. On the other hand, films with thicknessesgreater than several monolayers usually are adequateto simulate bulk-terminated surfaces.Therefore bulk optical constants may deliver indeed a reasonable description of the optical properties of uitra-thin layers.
Acknowledgements Financial support of this work by the Bundesministerium fir Bildung und Forschung (BMBF) under contract no. 0329172A9 and by rhe Fonds der Chemischenlndustrie is gratefully acknowiedged.
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