Microstructure and stoichiometry dependence of ion beam nitrides as a function of energy and...

Microstructure and stoichiometry dependence of ion beam nitrides as a function ofenergy and temperature: A comparative study between Si and SiGeOlof C. Hellman, Nicole Herbots, Olivier Vancauwenberghe, R. J. Culbertson, and W. J. Croft Citation: Journal of Vacuum Science & Technology A 10, 1631 (1992); doi: 10.1116/1.577761 View online: http://dx.doi.org/10.1116/1.577761 View Table of Contents: http://scitation.aip.org/content/avs/journal/jvsta/10/4?ver=pdfcov Published by the AVS: Science & Technology of Materials, Interfaces, and Processing Articles you may be interested in Ion dose, energy, and species dependencies of strain relaxation of SiGe buffer layers fabricated by ionimplantation technique J. Appl. Phys. 107, 103509 (2010); 10.1063/1.3374688 Strained silicon on SiGe: Temperature dependence of carrier effective masses J. Appl. Phys. 94, 5088 (2003); 10.1063/1.1609051 Ge composition and temperature dependence of the deposition of SiGe layers J. Appl. Phys. 75, 5382 (1994); 10.1063/1.356989 New SiGe dielectrics grown at room temperature by lowenergy ion beam oxidation and nitridation Appl. Phys. Lett. 59, 2031 (1991); 10.1063/1.106122 Growth temperature dependence of interfacial abruptness in Si/Ge heteroepitaxy studied by Ramanspectroscopy and medium energy ion scattering Appl. Phys. Lett. 54, 219 (1989); 10.1063/1.101014

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Microstructure and stoichiometry dependence of ion beam nitrides as a function of energy and temperature: A comparative study between Si and SiGe

Olaf C. Hellmana) and Nicole Herbots Department of Physics and Astronomy, Arizona State Uniuersity, Tempe, Arizona 85287

Olivier Vancauwenbergheb) Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139

R. J. Culbertson and W. J. Croft us. Army Materials Technology Laboratory, Watertowl!, Massachusetts 02172

(Received 12 November 1991; accepted 9 December 1991)

The microstructure and stoichiometry of nitrides formed by direct low-energy ion beam nitridation has been investigated as a function of ion energy and substrate temperature for Si( 100) and SiGe/Si( 100) films. Cross-sectional transmission electron microscopy, Rutherford backscattering spectroscopy combined with ion channeling and in situ x-ray photoelectron spectroscopy were used. It was established that a substrate temperature of 700 K produces a homogeneous amorphous nitride layer, whereas lower substrate temperatures decrease the incorporation of nitrogen in the film, while causing the formation of a nitrogen-poor amorphous layer beneath the nitride film. The N-to-Si or N-to-{Si + Oe) atomic ratio is found be close to 1.33 at 1 keY and decreases with ion energy. Effects due to chemically enhanced physical sputtering of germanium are observed.

I. INTRODUCTION Recent developments in the fabrication of SiGe-based het-erostructures on Si 1-3 have prompted interest in the forma-tion of SiGe dielectrics. As heteroepitaxial alloys are by nature metastable or unstable, their use necessitates pro-cessing methods with low thermal budget. Low energy ion processing has been used by several researchers as a low temperature technique to synthesize directly thin film in-sulating compounds of Si4-l9 and GaAs 20 based dielectrics. The formation of thin film nitride and oxide compounds of SiGe/Si has been recently demonstrated. 21 •22 In the case of oxidation of SiOe, purely thermal techniques have been shown to lead to rejection of Ge out of the SiGe alloy and growth of Si02• Ion beam oxidation (IBO) limits alloy decomposition due to limited diffusivities when an oxide is grown at room temperature. The kinetics of SiGe oxidation by both thermal and ion beam methods are the subject of continuing study.23-27 In this work, the kinetics of ion beam nitridation (IBN) and the properties of thin films of SiGe nitrides formed by IBN is investigated in order to identify and characterize the role of temperature and en-ergy of low energy IBN. The kinetics of IBN of pure Si21

are used as a basis for analyzing IBN of SiGe.

It EXPERIMENTAL PROCEDURE Single-crystal SiC 1(0) substrates were cleaned by a pro-

cedure (described in Ref. 21) based on a technique of Grunthaner et af., 28 using HF in ethyl alcohol. The initial surface exhibits only trace contamination from physisorbed oxygen and carbon as monitored by x-ray photoelectron spectroscopy (XPS).2i Such low temperature cleaning methods have been recently studied for application to molecular-beam epitaxy (MBE).29.30 Heteroepitaxial SiGe

films were grown on clean Si; this process is described elsewhere. 3 I Films were then transferred in UHV to a nitridation/XPS chamber. SiC 100) and SiGe/Si( 100) films were nitrided by direct exposure to N/ at energies ranging from 100 to 1000 eV, at temperatures ranging from 290 to 700 K. Chamber pressure rose from a base of I X 10 - 10 to 1 X 10 - 6 during nitridation. Ion current was measured at the sample bias power supply as the difference between the power supply current with the ion beam on and off. Ion currents were typi.cally 400 nA. The beam was rastered electrostatically to produce a beam spot of approximately 1 cm2 on the sample. Deposition was interrupted periodically to examine the film using in situ XPS. Samples were heated during IBN using a resistive heater located on the back side of the sample holder.

In situ XPS was conducted before and after IBN, and at regular time intervals during IBN. The MgKa emission was used with an electrostatic analyzer resolution of f:t..E/ E = 0.015 at takeoff angles of 36° and 45°. Films were ex-amined after IBN by XTEM using a JEOL 200-CX. Com-position of the films was measured by RBS combined with ion channeling using 2 MeV 4He + .

Two ion doses were investigated: a "low" dose of about 1 X 10\6 N 2+ /cm2 and a "high" dose of 3 X 1017 N 2+ /cm2.

For the low dose samples, nitridation was terminated at the point at which steady XPS emission was observed. The higher dose was chosen to be well in excess of the dose where steady XPS emission was observed, so that effects such as nitrogen incorporation deeper in the sample than the depth sensitivity of XPS and ion-beam-induced amor-phization could be investigated.

1631 J. 'lac. Sci. Technol. A 10(4), Jull Aug 1992 0734-2101/921041631-06$01.00 @1992 American Vacuum Society 1631


1632 Hellman et al.: Microstructure and stoichiometry dependence of ion beam nitrides 1632

250

200

.-- 150 '" '" 1:1

100 u

50

0 0 200 400 600 800 1000 1200

Ion Energy (e V)

FIG. 1. Amorphous layer thickness vs N2f ion energy for doses of 1 >< !O16

and 3X 1017 N/ fcm2• Ion projected range as calculated by TRIMCSR is shown for comparison.

III. RESULTS

A. TEM and STEM

Transmission electron microscopy (TEM) was used to determine the thickness of the ion beam synthesized nitride layer. After IBN, an amorphous layer is always observed on top of the Si ( 100) substrate or SiGe epitaxial film with a very sharp and uniform amorphous/crystalline interface. However, because TEM measurements are sensitive to structure rather than composition, and XPS depth sensi-tivity is limited, the composition of the film was established using scanning transmission electron microscopy (STEM).

When the low dose was used, the thickness of the amor-phous layer was found to be well beyond the ion projected range rp' and also beyond the end-of-range interstitials cre-ated by atomic collisions as calculated by TRIMCSR. 32 For

X-ray peak heights (Arbitrary Units)

lal

J. Vac. Sci. Technol. A, Vol. 10, No.4, JullAug 1992

ion energies of 200 eV (rp = 11 A) and 350 eV (rp = 16 A), the measured thickness was 40 A. The thickness de-creased to 26 A at 10 eV (rp = 6 A).

At high dose, however, the amorphous layer was found to grow thicker at lower ion energies. At 200 eV, for ex-ample, the amorphous layer thickness was 200 A rather than 40 A, even though the XPS emission remained con-stant (Fig. 1). High-dose samples were characterized by STEM. Because x-ray emission in STEM is affected by fluorescence and absorption probabilities, sample thick-ness, and surface contamination of the TEM sample as its upper and lower surfaces, absolute concentrations cannot be measured. Nevertheless, the nitrogen x-ray peak height is nearly proportional to the nitrogen concentration through the film. Figure 2 shows STEM results for IBN at 200 and 500 eV in the high-dose case. Bright-field STEM micrographs are shown with corresponding x-ray peak heights at selected spots through the films. For the 500 eV sample, the decay of the nitrogen concentration is found to correspond closely to the location of the microstructural interface, suggesting the formation of an amorphous ni-tride phase on a nitrogen-free crystalline phase. However, for the 200 eV sample, the nitrogen signal has decayed at a 100 A depth, suggesting that the substrate underlying the nitride is amorphized without significant incorporation of nitrogen.

The thickness of the amorphous layer was not found to be a function of deposition temperature. Low-dose samples prepared with 200 eV ions at 290,390,490,590, and 690 K all exhibited thickness 38 ± 2 A. However, differences were observed in the microstructure of the crystalline substrate immediately beneath the amorphous layer. Figure 3 shows bright-field electron micrographs of films deposited at these temperatures, with the sample aligned such that (110) is

Epoxy

Amorphous Layer

Substrate

Epoxy

Amorphous Layer

Substrate

FIG. 2. STEM micrographs and corre-sponding x-ray emission intensities for N, Si, and 0 through nitride films synthesized at Ca) 200 eV and (b) 500 eV.



lal

(b)

leI

FIG. 3. Bright-field TEM micrographs of nitride films synthesized at Ca) 690 K, (b) 490 K, and (c) 290 K. Alignment of all samples produces diffraction pattern shown in (d).

J. Vac. Sci. Techno!. A, Vol. 10, No.4, JuliAug 1992

tilted about 2° away from the electron beam. This align-ment is close to the strong diffraction condition which oc-curs when the electron beam is parallel to (110), while keeping the electron beam paranel to the sample surface. Precise alignment is made so that the electron beam dif-fraction pattern matches that shown in Fig. 3(d) exactly. Strain in the substrate due to defects will thus align the lattice so as to diffract strongly, producing contrast.

In Fig. 3(a), the structure grown at 690 K sample ex-hibits strained regions immediately beneath the amorphous/crystalline interface. These are interpreted as dislocation loops formed by the coalescence of interstitial point defects created by atomic collisions. At 490 K these regions produce less contrast, and at 290 K no contrast is observed. Such temperature dependence of the occurrence of buried damage has been observed previously in direct ion beam epitaxy33 and in lBO,18 and has been modeled theoretically.34 The appearance of dislocation loops at higher temperatures can be accounted for by the thermal activation of the diffusion and coalescence of these inter-stitials. At lower temperatures, the defects remain in the substrate, but do not coalesce into extended defects large enough to produce such contrast by TEM. The presence of defects in the region beneath the amorphous layer suggests that the mechanism of amorphization below the nitride layer is the collection of defects to the point of lattice col-lapse, similar to the mechanism proposed for high energy ion beam amorphization and ion-induced solid phase epitaxy. 35 We speculate that a transition temperature would determine the growth Of elimination of the amOf-phization layer in a series of samples implanted at a similar range of temperatures at the higher ion dose. Higher tem-peratures would suppress the growth of an amorphous layer, favoring instead the growth of extended defects, as in the low-dose case.

Only a small difference was observed between amor-phous layer thickness for IBN of SiO.89GeO.11 alloys com-pared to IBN of Si. For 290 K IBN using a lower dose of 200 e V ions, thickness of 44 A for the amorphous layers were measured. This is slightly higher than observed for IBN of Si. The conclusion drawn from this observation is that the addition of Ge to Si increases the rate of amor-phization of the substrate. This observation is in agreement with previous experimental studies of ion amorphization of Si and Ge. 36,37

B.XPS XPS has been used to identify chemical states of Si and

Ge in nitride films using the Si 2p (BE = 99.6 eV) and Ge 3d (BE = 29.7 eV) emissions. A chemical shift is observed in the energy of these signals when these elements are bonded to nitrogen. The ratio of the integrated areas of the shifted and unshifted peaks is taken to be proportional to the ratio of nitrided Si or Ge. Further, the integrated area of the shifted and unshifted peaks combined is taken to be proportional to the absolute concentration of that element near the surface. In this way, the fraction of the initial film which is nitrided can be calculated, and changes in the Si-to-Ge ratio in SiGe alloy films due to preferential sput-



] .:: 100 :.c [/l

90 Q.,

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.. J", II' !., ... [.""",.".".J." .. ,."., .. "."j" OJ 70 B ,··· .. · .. • .. · .... , .. ·t .... ·• .... · ...... · .. • .. ·; .. ,· .... , ...... · ...... +-0

Ef Po 60

N

! ········ .. t· , Vi J< 50

O. 200 400 600 800 1000 1200 Ion Energy

FIG. 4, Film stoichiometry vs ion energy, Percent of Si 2p and Ge 3d photoelectron peaks chemically shifted due to bonding with nitrogen.

tering can be monitored. XPS data (Fig. 4) shows that the film stoichiometry, expressed as the fraction of the Si and Ge photoelectron peaks which exhibit a chemical shift, is close to that of (SiGe)3N4 at higher ion energies, but be-comes nitrogen deficient at lower ion energies, as has been observed for the IBN of silicon.21 Figure 5 shows the evo-lution of the Si 2p and Ge 3d photoelectron peaks during IBN at 200 eV of a SiO.89GeO.11 alloy. More of the Si pho-toelectron peak is seen to be chemically shifted, as com-pared to the Ge peak, indicating a preferential bonding of nitrogen to Si. Secondly, the integrated area of the total Ge peak decreases with time. This indicates that depletion of Ge occurs at the surface due to the ion beam. This effect cannot be attributed to attenuation of the Ge 3d photoelec-tron signal by the nitrogen incorporated in the film, be-cause the Si 2p photoelectron, which is close in energy to the Ge 3d photoelectron, is not found to decrease. The depletion ofGe is a function of ion energy (Fig. 6). At low energies, depletion of Ge is low, while it increases at higher ion energies.

Depletion of Ge during IBN of SiGe alloys is attributed to a preferential sputtering effect. Following the definition of Betz and Wehner,38 the mechanism is preferential sput-tering could be (1) differences in physical sputtering, due to the mass and binding energy differences between Si and Ge during atomic collisions, (2) chemically enhanced physical sputtering, in which the chemical environment of Si and Ge surface atoms during IBN changes the relative binding en-ergy of these species, thus increasing the sputtering yield of

400 600 800 1000 1200 Ion Energy

FIG, 6. Ratio of Ge 3d signal integrated area after IBN to that before IBN as a function of energy.

one over the other, or (3) chemical sputtering, in which the ion beam catalyzes the synthesis of a volatile molecule at the surface, which then desorbs by thermal activation.

Physical sputtering and chemical sputtering cannot ac-count for the energy dependence of the sputtering effects observed. Physical sputtering can be calculated by TRIMCSR. 25 A similar simulation has been used recently for sputtering studies of SiGe alloys using high-energy oxygen ions with good agreement with experimental observations.39 TRIMCSR does not predict preferential sput-tering for IBN of SiGe alloys over the range of energies 100-1000 e V, using the lattice binding energies of 2 e V and surface binding energies of 1.8 eV for Ge and 2.8 eV for silicon.40,41

A chemical sputtering effect can be ruled out because the observed depletion of germanium is increasing contin-uously past 200 eV ion energy, which is well beyond the energy which would be required to catalyze the formation of a volatile species. Preferential sputtering of Ge due to a chemically enhanced physical sputtering, however, is ex-pected to increase with increasing ion energy. The mecha-nism responsible for Ge depletion must therefore be a re-duction in Ge surface binding energy when Ge is nitrided. Such an interpretation is also supported by our previous results showing that IBN of pure Ge does not result in the formation of a Ge nitride. 21 ,22 In this case, chemically en-hanced physical sputtering of Ge removes the surface ni-tride layers at a rate higher than the rate of nitrogen in-

Ge

Ge3d l

FIG. 5, Evolution during nitridation of Si 2p and Ge 3d photoelectron signals for 200 eV nitridation of SiQ,89GeO.11 alloy,

I 36

I 32

i 28

I 24

Binding Energy (eV)

J. Vac. Sci. Techno!. A, Vol. 10, No.4, Jul/Aug 1992



20 ! N 16 E u

---'" 12 a § ...

J '" Z 8

+ .......................... + ........................... +.", ........................... , .......................... r'·t ...... ·· .. -i-'

t j ................................. , ...................... ", ..... -+ ....................................................................................... al-

V)

0

>< 4 .................................. , ............................................................... + ........................... + ......................... ....,

0 0 100 200 300 400 500

Temperature (0C)

FIG. 7. Areal concentration of N as measured by RBS as a function of temperature during nitridation.

corporation, even at energies as low as 100 eV.

C.RBS Rutherford backscattering spectroscopy (RBS) in com-

bination with ion channeling was used both to analyze the composition of the SiGe alloy films and to quantify the incorporation of nitrogen of the films. When combined with stoichiometry measurements by XPS, nitrogen incor-poration as measured by RBS can be used to predict nitride layer thicknesses, which can then be compared to TEM measurements.

The nitrogen incorporation for IBN of Si at a fixed dose increases with temperature from RT to 690 K (Fig. 7), demonstrating that thermally activated kinetics are par-tially responsible for the chemical reaction of the nitrogen ion beam with the substrate. On the other hand, given the very low equilibrium diffusivity of N in Si42 or Si3N4,43 and the very high temperatures (T> 1100 K) required for a thermally driven nitridation of silicon,44 the temperature used during IBN is far too low for the nitridation reaction to occur by a purely thermal kinetic path. The kinetics of IBN are neither controlled by atomic collisions only, as this would be temperature independent. Thus, atomic col-lisions and thermal effects both participate in the IBN ni-tridation mechanism. Recent theoretical modeling taking into account both effects are shown to be consistent with these observations.45

The measured values for amorphous layer thickness by TEM correspond to thicknesses predicted by the combina-tion of RBS and XPS results. A 40 A film of SiN with a density of 2.9 gl em 3 (a compromise between densities of Si and Si3N4 ) would incorporate 1.6 X 1016 N atoms, which is the nitrogen content measured by RES of the film depos-ited at 690 K. The nitrogen content of the films grown at lower temperatures decreases, although the amorphous layer thickness remains constant and the XPS emission remains constant. An amorphous layer beneath the nitride is thus formed by ion beam amorphization, while the ni-trogen profile in the films synthesized at lower tempera-tures is determined by diffusion. That ion-induced amor-

J. Vac. Sci. Technol. A. Vol. 10, No.4, Jul/Aug 1992

phization does not occur at the higher temperatures is consistent with the temperature dependencies of ion-induced solid-phase epitaxy.36,37

Compared to IBN of Si, the nitrogen incorporation was found to decrease significantly when SiGe is nitrided: at 20 eV and 290 K, IBN of SiGe at the lower dose resulted in half the nitrogen incorporation observed for IBN of Si. This difference implies that a chemical effect reduces the amount of nitrogen incorporated in the film, and that a larger fraction of the film thickness is a nitrogen-poor amorphous SiGe phase, as compared to IBN of pure Si.

The heat of formation of Ge3N4 has been determined46

to be - 15.6 kcal/mol, compared to - 170 kcallmol for Si3N4.44 This difference can be invoked to explain the lower incorporation of nitrogen in SiGe films: given a lower driv-ing force to bond with the SiGe film, a free nitrogen atom is more likely to migrate to the surface, bond with another nitrogen atom, and desorb as a nitrogen molecule. In other words, chemical sputtering of nitrogen may be enhanced in SiGe alloys. This difference can also explain the trend for silicon to be preferentially nitrided. as shown in the XPS peak shifts.

IV. CONCLUSIONS In summary, the microstructure as a function of tem-

perature ofIBN of Si(100) has been investigated by TEM, and a comparative study of IBN of Si and SiGe has been concluded, High-dose IBN at room temperature leads to the buildup of an amorphous phase beneath the nitride. A similar trend is observed for SiGe. The growth of the amor-phous phase decreases with increasing temperature, and vanishes at 690 K. These are also the conditions at which the nitride phase is found to be mostly stoichiometric and also strain free. 21 However, buried damage is then observed below the interface. These results suggest that to optimize dielectric formation by IBN for such materials, a low-dose nitridation near 690 K should be used, combined with se-quential Si or SiGe deposition to overcome thickness lim-itations.

ACKNOWLEDGMENTS The authors would like to thank W. Kosik of the U.S.

Amry Materials Technology Laboratory for her expert as-sistance in performing the RBS and ion channeling analy-sis.

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