Raman monitoring of wide bandgap MBE growth

ELSEVIER Applied Surface Science 123/124 (1998) 276 282

applied sur face sc ience

Raman monitoring of wide bandgap MBE growth

D i e t r i c h R.T. Z a h n *

Halbleiterphysik, TU Chemnit:, D-09107 Chemnit=, Germany

Abstract

The combination of Raman spectroscopy and molecular beam epitaxial (MBE) growth can supply valuable additional information on interface and layer formation. Here the growth of CdS on InP(100), that of undoped and nitrogen doped ZnSe on GaAs(100) as well as the nitridation of GaAs(100) serve as examples for the usefulness of this in situ and on line

technique. In particular, evidence for a reaction in the initial phase of CdS/InP(100) interface formation and the transformation from the cubic to the hexagonal phase of CdS at around 200 nm layer thickness is provided. In the case of ZnSe growth on GaAs(100) the incorporation of nitrogen in the layer results in a deterioration of the ZnSe layer crystallinity together with the appearance of compressive strain in the GaAs substrate. The latter is assigned to nitrogen diffusion into the substrate which is further supported by a nitridation experiment of GaAs which also leads to the formation of a thin layer of GaN. Moreover, the underlying substrate is heavily disturbed as can be judged from the detection of As clusters. © 1998 Elsevier Science B.V.

1. Introduction

In recent years Raman spectroscopy has been further developed to a standard that it can now be applied in situ and on l ine during the growth of semiconductor layers by MBE. The set-up required for this so-called Raman monitoring combines an ultra-high vacuum (UHV) chamber for MBE and a multichannel Raman spectrometer which allows for fast recording of the Raman spectra [1], Another prerequisite of this method is the utilization of resonance enhancement of Raman scattering in order to increase the sensitivity for the detection of ultra-thin layers in the nanometer range. For the wide bandgap materials studied so far, i.e. CdTe, CdS, ZnSSe, and GaN, this resonance enhancement is usually achieved by excitation with a photon energy in the vicinity of the E 0 gap. Here the temperature induced bandgap

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shift has to be taken into account since growth is usually performed at elevated temperature. The effect of temperature induced resonant Raman scattering was thoroughly studied for ZnSSe layers on GaAs [2]. There it was shown that the scattering intensity can be increased by approximately two orders of magnitude compared to a far off-resonance condition and that one can obtain further valuable information on e.g. the temperature dependence of the energy gap. By carefully adjusting the growth temperature and the excitation photon energy it is thus possible to obtain sufficient scattering intensi- ties of very thin wide bandgap semiconductor layers with good signal-to-noise ratio even at very high temperatures as demonstrated for the growth of GaN at above 600°C [3].

A considerable amount of the work concerning Raman monitoring of wide bandgap semiconductor growth has already been published and therefore only some of the main results are briefly reviewed in the following. The deposition of CdTe on the lattice

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D.R. 7". Zahn / Applied Surface Science 123 / 124 (I 998) 276-282 277

matched substrate InSb is an example for a highly reactive interface [4]. It was confirmed that the deposition of CdTe at elevated temperature (300°C) results in the formation of In2Te 3 and liberated Sb rather than CdTe growth. Even at room temperature deposition Te was found to be incorporated into the InSb substrate leading to an increased carrier concentration. In comparison to the CdTe/InSb system the growth of ZnSSe on GaAs is extremely unreactive so that information deduced mainly concerns the growing layer [5,6]. In particular, the composition of the ternary compounds can be derived from the longitu- dinal optical (LO) and 2LO features in the resonant Raman spectra with an accuracy of better than 1%. Finally the deposition of GaN on strongly lattice mismatched substrates like GaAs and Si revealed a severe influence of the layer growth on the substrate, namely the evolution of compressive strain near the interface [3].

Here the capabilities of Raman monitoring will be demonstrated using three recent studies of CdS on InP(100), the comparison of undoped and nitrogen doped ZnSe on GaAs(100), and the nitridation of GaAs. CdS epitaxial growth was previously studied by in situ but not on line Raman spectroscopy on the (110) surface of lnP [7]. The results obtained now for the growth on the (100) surface of InP by Raman monitoring substantiate the detection of a interface reaction and the transition from the substrate stabilized cubic phase to the naturally preferred hexagonal phase at around 200 nm [8]. Considering the comparison of undoped and N doped ZnSe not only a deterioration of the ZnSe is detected as a result of nitrogen doping but this example also reveals another very important advantage of Raman monitoring, namely its capability to follow changes occurring in the near interface region of the substrate as growth progresses. This sensitivity for buried inter- faces is then further exploited in the GaAs nitridation experiment which demonstrates the formation of a thin GaN layer but also that of subsurface As clusters.

2. Experimental

MBE growth was performed in a specially de- signed UHV chamber (base pressure < I X l O - l °

mbar) which allows light scattered from the sample surface to be collected through a central viewport and focused onto the entrance slit of a Raman spectrometer (Dilor XY) with multichannel detection. The resolution varied between 3 and 5.5 cm ~ de- pending on the excitation photon energy used. De- tails of this experimental set-up can be found in Ref. [1]. Substrates used for CdS growth were (100) oriented InP wafers which were pretreated in ammo- nium sulfide and thereafter annealed in UHV at 330°C for 10 min leading to sulphur terminated surface. For ZnSe growth and nitridation micron thick homoepitaxial GaAs layers grown by MBE and thereafter capped with a thick layer of amorphous As were employed. A clean c(4 × 4) reconstructed surface was prepared by thermal desorption of the As cap. CdS and ZnSe were evaporated from Knudsen cells containing high purity polycrystalline com- pound source material while atomic nitrogen was provided by a if-plasma source (SVT RF 4.5). In some cases a quartz crystal microbalance in the vicinity of the substrate was used to monitor the deposition. Substrate temperatures during growth were 200°C and 260°C for CdS and ZnSe growth, respectively, while GaAs nitridation took place at 750°C. All these temperatures were derived from the temperature induced frequency shift of the substrate LO phonons. Above room temperature this shift follows very well a linear dependence and was measured to be - 2 . 0 and - 1 . 8 cm- I /100°C for InP and GaAs, respectively. An accuracy in the determi- nation of the frequency position of better than + 0.2 cm 1 can easily be achieved thus resulting in an error for the surface temperature of < _+ 10°C.

3. Results and discussion

3.1. CdS g r o w t h on InP( lO0)

The evolution of Raman spectra as function of CdS deposition time on InP(100) is displayed in Fig. 1. The first featureless spectrum corresponds to that of a clean substrate since the scattering geometry with the incident and scattered light polarized along a [010] direction of the substrate only allows F65hlich induced scattering by the CdS LO phonons to be observed. As can be clearly seen CdS LO scattering

278 D.R.T. Zahn /Applied Sur[ace Science 123/124 (1998) 276-282

intensity / counts mWls 1

0 . 2 -

0.1

0 , 0 - - \\

~0

500 / ~ \ 120 Raman Shift / cm 1 600 0 time I min

Fig. 1. Raman monitoring of CdS growth at 200°C on InP(100) using a photon energy of 2.34 eV.

appears near 300 cm ~ in the spectra after a few minutes of deposition while the onset of 2LO scattering near 600 cm ~ is somewhat delayed. The observation of both is induced by the resonant excitation using a photon energy of 2.34 eV and the 2 L O / L O ratio critically depends on the difference between bandgap and photon energy. Superimposed on the general intensity increase is a modulation of the phonon features which is a consequence of Fabry- Perot interference of incident and scattered light. This effect is observed for all cases for which a homogeneous layer with optically smooth surface and interface is formed. The intensity modulation can be simulated using the optical constants of layer and substrate material [9] and then provides values for the total thickness and the growth rate. The time scale in Fig. 1 can therefore be converted into a thickness scale. Here a layer thickness of 210 nm

was obtained at a growth rate of approximately 0.6 nm/min .

Closer inspection of the spectra of Fig. 1 recorded in the initial phase of deposition reveals the presence of a broad shoulder on the low energy side of the LO phonon feature. This shoulder has already previously been assigned to the formation of an In -S rich layer at the interface [7]. Its observation indicates that even though the InP(100) surface is S-terminated the sulphur provided by the molecular beam further reacts with the substrate to form this In -S rich layer prior to CdS epitaxia] growth. In agreement with the previous studies for CdS growth on InP( l l0) the formation of this thin reacted layer is no obstacle for good epitaxial growth.

A second issue addressed for CdS growth is the transition of the substrate stabilized cubic modification into the naturally preferred hexagonal modifica-

D.R.T. Zahn /Applied Surface Science 123/124 (1998) 276-282 279

tion which was reported to occur around 200 nm for growth on (110) oriented InP [8]. For this purpose layers with thicknesses much higher than 200 nm were grown and monitored by Raman as well as by photoluminescence (PL) spectroscopy using excitation at 3.81 eV. The latter has been attempted for the first time and provides important additional information as anticipated from previous investigations [10]. The Raman spectra of the hexagonal modification differ significantly from those of the cubic one only for high quality bulk crystals for which additional modes can be detected, whereas the spectra of thin films in particular under resonance condition show very similar LO and 2LO features. However, due to the difference in bandgap of approximately 100 meV [8] there is a change in the resonance condition which should be observable by a corresponding change in the 2 L O / L O intensity ratio. The PL spectra (not shown) consist of a single feature. Its maximum position is plotted in Fig. 2 as a function of CdS layer thickness together with the evolution of the 2 L O / L O ratio. In the logarithmic scale used here the thickness begins at 10 nm nominal deposi-

2.38

2.36 > ll) ~ 2 . 3 4

2.32 _..I o_ 2.30

2.28

2.26

0.6 0

0.5 0

0.4 0 '~ 0.3 n -

0.2 t,,o

0.1 c-

- 0.0

D 0 0

P L

Eex = 3 8 1 e V , e e "

I ,

%

,k ] 1

• Raman • • E =2.34eV

ex

J . . . . . . t

0 100

O N

1000

CdS Layer Thickness / nm

Fig. 2. Evolution of CdS PL energy and 2LO/LO ratio as a function of layer thickness.

tion since this offset was found be present between the deposition on the quartz crystal and the start of CdS growth on the sample. This is likely to be caused by the consumption of the first 10 nm nominal deposition for the interface reaction. Considering the PL energy there is a clear blue shift for very thin layers which is assigned to confinement. Then the energy is constant at about 2.31 eV up to almost 30 nm and thereafter decreases to a minimum at about 2.27 eV near 200 nm followed by a final increase to 2.34 eV. This behavior can be understood assuming that the CdS layer grows pseudomorphically on InP up to the critical thickness of about 20 nm and that then the strain is released by the formation of dislocations. The minimum value of 2.27 eV should therefore be close to the PL energy for relaxed CdS in the cubic phase while the following shift towards higher energy is indicative of the formation of a larger bandgap material, namely the hexagonal phase of CdS. The difference of almost 80 meV between the minimum and the final value should thus be close to the difference in the bandgaps of the two modifications at 200°C. Considering now the evolution of the 2 L O / L O ratio it can be seen that it first increases monotonically with thickness revealing that the 2LO scattering intensity exhibits its full strength only in bulk like crystals. Their is a minor indication of a change in slope near 30 nm, i.e. near the onset of relaxation via dislocations. The marked change, however, is apparent at 200 nm again where the transition to the hexagonal phase and the corresponding shift in bandgap energy cause a change in the resonance condition resulting in a downshift of the 2 L O / L O ratio.

3.2. Undoped and nitrogen doped ZnSe on GaAs(lO0)

The growth of ZnSe was in both cases performed under identical conditions except that additional atomic nitrogen was provided by the rf plasma source operating at a power of 400 W and a nitrogen flow of 2 sccm. The doped layers grown reveal p-type conductivity with carrier concentrations around 5 × 10 ~6 cm 3. The Raman spectra were recorded using a photon energy of 2.56 eV at the growth temperature of 260°C in a backscattering geometry with both polarizations of incident and scattered light parallel

280 D,R. 7", Z~thn /Applied Surface Science 123/124 ([998) 270-282

to a [011] direction of the substrate. This allows the scattering by the GaAs LO phonon also to be observed. A comparison of the evolution of the Raman

spectra for undoped and N doped ZnSe is displayed in Fig. 3. The time scale has already been converted into a thickness scale by simulating the obvious

(a) Intensity / counts mW i s -1

S j ,~'~'~ / ~ " ~ "

0.10

0.08

0.06

0.04

0.02

-q

GaAs LO ................................. ZnSe 2 LO

100 200 -- '-"- '------~- 300 400

Raman Shift / cm "l 0

50 200

5O Thickness !nm

(b) Intensity / counts mW i s -1

0.10

0.08

0.06

0.04 20g 5° 0.02

400 500 6~ :~ '~ "~=rO 50 Thickness i nm Raman Shift / cm "1

Fig, 3. Raman monitoring of undoped (a) and nitrogen doped (b) ZnSe on GaAs at 260°C using a photon energy of 2,54 eV

D.R.T. Zahn /Applied Surjace Science 123/124 (1998) 276 282 281

intensity modulations. The three features visible correspond to scattering by the GaAs LO, the ZnSe LO, and ZnSe 2LO at approximately 285, 245, and 490 cm l, respectively. There is a drastic difference since the 2LO scattering is hardly visible in the case of doped ZnSe whereas it is fairly strong in the undoped case. As the 2LO scattering is generally very sensitive to the crystal perfection and is reduced in the presence of lattice disturbance the weak 2LO scattering already indicates that a lattice deterioration goes along with the nitrogen incorporation. This is further supported by directly comparing two Raman spectra in the LO range at the same layer thickness of 200 nm as shown in Fig. 4. Here it can clearly be seen that the nitrogen incorporation leads an increase in the full width at half maximum of the ZnSe LO from 13.7 to 20.2 c m - i. While the effect of the free carriers on the LO lineshape at the concentration measured is almost negligible [11], the broadening can almost entirely be attributed to lattice disorder as a consequence of nitrogen incorporation.

The presence of nitrogen during growth does not only affect the layer crystallinity but also induces severe changes in the near interface region of the substrate, i.e. within the Raman information depth of approximately 40 nm. This is clearly illustrated by the evolution of the LO frequency position with layer thickness as shown in Fig. 5. The LO frequency remains almost constant for the growth of undoped ZnSe with a minor tendency towards lower frequency in agreement with a possible tensile strain in the substrate as a result of the lattice mismatch.

O l O i ' i i ' , t ' i . i ' i ' i

Z n S e LO T 0 0 9

E 0 08 d

E 0 07

0 0 6

005

{B 0 04

0 03 E

D~ 002 , t , t , t , t , t , t , t , I , t ,

125 150 175 200 225 250 275 300 325 350 375

Raman Shif t / cm -1

Fig. 4. Comparison of the Raman spectra of undoped and nitrogen doped ZnSe on GaAs in the ILO range at the same thickness of 200 nm.

i i i i i i i

288 8 o o ° o o o

o o o ° o 286 5 0 ° o o 0 o o

o o o Oo o o o o

[ o ZnSe:N 1 .~_ 2 8 5 9 J • Z n S e u n d o p e d +

2 8 5 6

~. 2858 • 2 ' ." " _ ~ • .. - . " . ,

2 8 5 0 •

2 8 4 7 I , I 0 , I , I , I , I , I 0 5 100 150 200 250 300

T h i c k n e s s / n m

Fig. 5. Evolution of GaAs LO phonon position upon ZnSe deposition for undoped and nitrogen doped ZnSe.

Growth of N doped ZnSe, on the other hand, induces a shift towards higher frequency indicating a compressive strain in the near GaAs interface region. This is most likely due to the incorporation of nitrogen in the GaAs lattice. The average strain within the information depth amounts to approximately 500 MPa. A qualitatively very similar behavior was also found for the growth of GaN on GaAs [3].

3.3. GaAs nitridation

The nitridation was performed on decapped GaAs MBE layers at a temperature of 750°C. The Raman spectra were recorded in the same geometry as the ZnSe growth but using a photon energy of 3.05 eV which is close to the bandgap of GaN at the growth temperature and has a very low penetration depth of 18 nm into GaAs. Fig. 6 shows a few selected spectra from various nitridation stages. While there is hardly any change after one hour of nitridation and the spectrum is dominated by the GaAs LO scattering, new features evolve below the GaAs LO peak and just above 700 cm i. The latter can be unam- biguously be assigned to scattering from a thin GaN layer formed by the nitridation process [3]. its thickness is approximately 20 nm as can be judged from a comparison with Raman spectra recorded during GaN growth [3]. Similar to previous GaN growth experi- ments and nitrogen doping of ZnSe the GaAs LO phonon frequency shifts upwards indicating the lor- mation of strain up to 1 GPa in GaAs upon nitridation. The features occurring below the GaAs LO

282 D.R.T. Zahn /Applied Surface Science 123/124 (1998) 276 282

> . ,

. m

t - ~9 ¢ -

e-

ro O0

i

0

/ hv = 3 . 0 5 e V

6h

5h

4h

3h

2h

200 400 600 800 1000 1200

Raman Shif t / cm -1

Fig. 6. Selected spectra recorded at 750°C using a photon energy of 3.05 eV from various stages of the nitridation of GaAs.

stand the interface formation. The particular strength of the method lies in its ability to detect signals from the growing overlayer and of the underlying substrate at the same time. It is therefore possible to monitor e.g. the evolution of strain in the substrate as

demonstrated for nitrogen doped ZnSe as well as for the GaAs nitridation. Furthermore, it was shown that the method under the resonance conditions used is extremely sensitive to changes of crystallinity a n d / o r electronic properties and that the in situ monitored photoluminescence is also extremely useful for the understanding of semiconductor growth.

Acknowledgements

I am very grateful to D. Drews, A. Schneider and N. Hartmann for their experimental work, to K. Horn and D.A. Evans for their support of the C d S / I n P experiment and D. Westwood for the supply of As capped GaAs material and finally the DFG for the financial support under the grants Za 146 / 2 and Za146/7-1.

peak were identified by additional Raman spectra taken post-growth using a photon energy of 1.92 eV which is close to resonance with arsenic which was

assumed to be a probable cause for the scattering intensity 200-250 cm i. Indeed these spectra clearly

reveal the presence of crystalline As clusters as can be judged from the detection of the Eg and Alg modes of crystalline As. These As clusters have to be located below the surface GaN layer since any surface As would have been evaporated from the surface at the high growth temperature. In addition,

it was found that annealing in UHV or N 2 atmo- sphere does not have any significant effect on the Raman spectra of GaAs.

4. Summary

The examples given above reveal that Raman spectroscopy applied in situ and on line during the growth of semiconductor layers can help to under-

References

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[4] D. Drews, J. Sahm, W. Richter, D.R.T. Zahn, J. Appl. Phys. 78 (1995) 4060.

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Srama, U. Rossow, A. Thomas, K. Horn. W. Richter, Appl. Surf. Sci. 56-58 (I992) 684.

[8] D.R.T. Zahn, G. Kudlek, U. Rossow~ A. Hoffmann, I. Broser, W. Richter, Adv. Mater. Opt. Elect. 3 (1994) 11.

[9] V. Wagner, W. Richter, J. Geurts, D. Drews, D.R.T. Zahn, J. Raman Spectrosc. 27 (1996) 265.

[10] A. Kelnberger, G. Kudlek. U. Rossow, M. Ktipp, N. Esser, A. Hoffmann. W. Richter, D.R.T. Zahn, in: Proc. ICFSI-4 Jiilich, World Scientific, 1993, p. 696.

[11] D.J. Olego, J. Petruzzello, T. Marshall, D. Cammack, J. Cryst. Growth 117 (1992) 324.

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Raman monitoring of wide bandgap MBE growth