- Email: [email protected]
Amorphous inclusions during Ge and GeSn epitaxial growth via chemical vapor deposition
Thin Solid Films 590 (2015) 163–169
Contents lists available at ScienceDirect
Thin Solid Films journal homepage: www.elsevier.com/locate/tsf
Amorphous inclusions during Ge and GeSn epitaxial growth via chemical vapor deposition F. Gencarelli a,b,⁎, Y. Shimura a,c,d, A. Kumar a,c, B. Vincent a, A. Moussa a, D. Vanhaeren a, O. Richard a, H. Bender a, W. Vandervorst a,c, M. Caymax a, R. Loo a, M. Heyns a,b a
imec, Kapeldreef 75, 3001 Leuven, Belgium Dept. of Metallurgy and Materials Engineering, KU Leuven, B-3001 Leuven, Belgium Nuclear and Radiation Physics Section, KU Leuven, B-3001 Leuven, Belgium d FWO Pegasus Marie Curie Fellow b c
a r t i c l e
i n f o
Article history: Received 10 September 2014 Received in revised form 30 July 2015 Accepted 30 July 2015 Available online 1 August 2015 Keywords: Germanium tin Amorphous Islands Epitaxial breakdown Low temperature chemical vapor deposition Digermane
a b s t r a c t In this work, we discuss the characteristics of particular island-type features with an amorphous core that are developed during the low temperature epitaxial growth of Ge and GeSn layers by means of chemical vapor deposition with Ge2H6. Although further investigations are needed to unambiguously identify the origin of these features, we suggest that they are originated by the formation of clusters of H and/or contaminants atoms during growth. These would initially cause the formation of pits with crystalline rough facets over them, resulting in ring-shaped islands. Then, when an excess surface energy is overcome, an amorphous phase would nucleate inside the pits and fill them. Reducing the pressure and/or increasing the growth temperature can be effective ways to prevent the formation of these features, likely due to a reduction of the surface passivation from H and/or contaminant atoms. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Ge and GeSn alloys are interesting group IV semiconductors with potentially superior transport [1,2] and optical [3,4] properties as compared to conventional Si. Low temperature growth processes are often required in order to fabricate Ge or GeSn heterostructures without undesired strain relaxations, elemental intermixing, dopant diffusion or Sn precipitation [5]. However, low-temperature growth is known to have a series of possible undesired consequences, such as the roughening of the surface [6], the formation of defects or the inclusion of localized or extended amorphous regions [6–8] (often referred to as epitaxial breakdown). Using low temperature chemical vapor deposition (CVD) we demonstrated the epitaxial growth of good quality, dislocation-free thin GeSn layers [9,10] using a high-order, more reactive Ge precursor as compared to standard GeH4 (namely Ge2H6) to achieve sufficiently high growth rates. However, we also observed that when the thickness of these layers increases, their surface is increasingly covered by islands, which we initially attributed to strain relaxation [7]. In this work, we investigate in more details the characteristics of these islands and we discuss the presence of an amorphous core inside them. Since the ⁎ Corresponding author at: imec, Kapeldreef 75, 3001 Leuven, Belgium. E-mail address: [email protected] (F. Gencarelli).
http://dx.doi.org/10.1016/j.tsf.2015.07.076 0040-6090/© 2015 Elsevier B.V. All rights reserved.
presence of these particular features is expected to complicate the fabrication of optical or electrical devices containing Ge or GeSn heterostructures and to have a negative impact on their performances, their formation mechanism is investigated in order to understand and possibly prevent their appearance. 2. Experimental details The analyzed Ge and GeSn layers are epitaxially grown on Ge(001) or on 1 μm Ge-buffered [11] Si(001) substrates in a 200 mm ASMEpsilonTM-like CVD reactor, employing a growth temperature of 320 °C and a growth pressure of 101 kPa (unless otherwise specified). Ge2H6 (diluted in H2) and SnCl4 are used as the Ge and Sn precursors, respectively, with either N2 or H2 as the carrier gas [9]. In order to remove the native oxides, a pre-Epi cleaning of the wafers is performed, consisting of a HF (2%) dip for 90 s followed by a rinsing step in deionized water and a Marangoni Dry step. In addition, the wafers received a standard in-situ pre-Epi bake under H2 for 10 min, at 650 °C and 5333 Pa. The surface morphology of the grown layers is inspected by atomic force microscopy (AFM) measurements, which are executed in tapping mode on an Icon PT tool coupled with a Nanoscope V controller (Bruker), using an AC160TS Olympus probe. The layer crystallinity is characterized using plan-view (PV) and cross-section (X) transmission
164
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
electron microscopy (TEM) together with a fast Fourier transform (FFT) analysis. Atom probe tomography (APT) and energy dispersive X-ray spectroscopy (EDS) are employed to analyze the distribution of the Sn atoms inside the GeSn islands. The APT analysis was done with the Laser Assisted Wide Angle Tomographic Atom Probe (LAWATAP) tool from Cameca using laser pulsing (wavelength 515 nm, 400 fs pulse duration) at 15 K sample temperature. The raw data were reconstructed using TAP3D data software from Cameca using the ‘Voltage curve’ reconstruction method [12]. Tips suitable for APT analysis were prepared by the lift out method and sub-sequential annular Focused Ion Beam (FIB) milling, on a FEI NOVA-600 dual beam tool. The TEM, EDS and high angle annular dark field scanning transmission electron microscopy (HAADF-STEM) analyses are performed on a Tecnai F30 ST tool from FEI using a FEG electron source operated at 300 kV. The X-TEM specimens are prepared by FIB milling (Strata, FEI). A CVD glass, a sputtered Al and a sputtered Pt layers are deposited beforehand on the area of interest. The PV-TEM specimens are prepared using a precision ion polishing system (PIPS, Gatan). The spreading resistance of the GeSn layers is measured in high vacuum (1.3 × 10−3 Pa) by scanning spreading resistance microscopy (SSRM) using a Bruker E-scope with a 50 mV sample bias and a deflection setpoint of 0.5 V. Finally, secondary ion mass spectrometry (SIMS) and total reflection X-ray fluorescence (TXRF) are used to address the incorporation of contaminant atoms inside the GeSn layers. SIMS measurements were performed using an Atomika 4500 tool, with a Cs + primary beam having an incidence angle of 45° and an impact energy of 5 keV. TXRF measurements are performed with an Atomika-FEI TXRF8300 25031 tool equipped with a W X-ray tube, using 50 kV and 55 mA as the X-ray tube voltage and current, respectively, and a measurement angle of 2.3 mrad. It is worth mentioning that the results of all these analysis are consistent when performed for samples of different thickness (in the range 30–1000 nm) and composition (in the range 0–12 at.% Sn). The pictures discussed in this work are purely selected according to their quality in terms of contrast and brightness, for the sake of clarity. The structural details of the corresponding samples are summarized in Table 1.
3. Results and discussion As previously reported [7], particular island-type features are formed on the surface of the growing Ge or GeSn layers during low temperature CVD experiments with Ge2H6. These islands, shown in the top view 2D AFM images of Fig. 1(a) for Ge (sample A1, top panel) and GeSn (sample B1, bottom panel), respectively, have different features, shapes and sizes, as it will be discussed subsequently. In the case of pure Ge growth, the X-TEM inspection (Fig. 1(b), top panel) reveals that inside each crystalline island there is an inverted pyramidal defect containing amorphous material. These defects have their symmetry axis along the growth direction and appear to be delimited by {112} facets, which form an angle of 35° with the (001) growing surface. As it is shown in Fig. 1(b) (bottom
panel), the GeSn islands contain similar defects, but apparently delimited by {221} facets forming an angle of 71° with the (001) plane. Due to the poor contrast between the crystalline and amorphous regions of the GeSn sample B1 in the X-TEM image in Fig. 1 (b) (bottom panel), PV-TEM images with increasing magnification factors are also shown for this sample in Fig. 2(a), (b) and (c). We observe that the top central region of each island has always a square shape, with rounded corners and sides along [010] directions, i.e. 45° rotated as compared to the external island sides aligned along [110] directions. This suggests that the inverted pyramidal defects inside the GeSn islands are actually delimited by {041} crystalline facets, as it is schematically shown in Fig. 2(e). Such a 45° rotation is not observed in the AFM image of the Ge layers (upper panel of Fig. 1(a)). We also observe that the stressinduced bend contours patterns present inside each GeSn island do not continue in its central region (Fig. 2(a) or (b)) and that diffuse rings appear in the FFT image (Fig. 2(d)) obtained from the highest resolution PV-TEM picture (in Fig. 2(c)). These findings indicate the presence of amorphous material inside the crystalline inverted pyramidal defects, also clearly visible in Fig. 2(c). Fig. 3 shows line scans across different islands seen in the top panel of Fig. 1(a) (corresponding to a 70 nm thick Ge layer), but plotted on the same x-axis. By overlapping the line scan profiles in this way, it is possible to observe the evolution of the islands for increasing deposition times. At their beginning stage, they appear as elevated “squared rings” with a dip in the middle (i.e. a pit). Subsequently, at some critical thickness (~ 5 nm), the excess surface energy associated with the pit walls is likely too large and the pit gradually fills in. However, the newly nucleated phase is amorphous and has a higher growth rate than the surrounding crystalline material in the ring, which continues its vertical and lateral growth. APT analysis is also used in order to characterize the amorphous inverted pyramids, which can be localized in the APT tip via the method detailed in Ref. 13. The 3-D reconstruction of a tip obtained from the GeSn sample C1 is shown in Fig. 4(a). For a clear visualization, Sn atoms are not visualized and only 20% of the Ge atoms are shown. We notice that during the final cleaning of the tip at low energies some Pt atoms (coming from a cap previously deposited on the GeSn layer in order to protect it from major ion beam damages) preferentially diffuse towards the amorphous region. This artifact is beneficial to easily visualize the inverted pyramidal defect, whose facets are found to form 71° with the (001) plane, in good agreement with X-TEM results. A 1-D composition profile – extracted from a region across the pyramid walls with the depth aligned along the cross-section of the tip (indicated by a black box in Fig. 4(a)) – is shown in Fig. 4(b). The Sn concentration inside the amorphous pyramid is found to be lower (8–10 at.%) as compared to that in the surrounding crystalline matrix (15 at.%). This variation in the Sn concentration is reflected by a complementary variation in the Ge concentration. It should be mentioned that these APT measurements were performed at a lower pulse fraction than the optimal one necessary to accurately quantify the absolute composition of the GeSn layers [13], which is then slightly overestimated. A lower Sn content in the amorphous GeSn region as compared to the crystalline area is also determined by EDS measurements of the Sn
Table 1 Samples discussed in this work. Sample ID
Epilayer thickness [nm]
Sn content [%]
Carrier gas
Substrate
Comment
A1 A2 A3 B1 B2 C1 D1 E1 F1
70 70 70 144 542 1000 27 295 67
0 0 0 6.7 6.1 12.7 6.4 0 0
N2 N2 H2 N2 N2 N2 N2 N2 N2
1 μm Ge/Si(001) 1 μm Ge/Si(001) 1 μm Ge/Si(001) 1 μm Ge/Si(001) 350 nm Ge/Si(001) 1 μm Ge/Si(001) Ge(001) 1 μm Ge/Si(001) 1 μm Ge/Si(001)
Reference No pre-EPI bake H2 carrier – – – Fully strained 10 Torr 400 °C
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
165
Fig. 1. (a) Top view 2D AFM and (b) X-TEM images of samples A1 (top) and B1 (bottom), indicating the presence of particular island-like features with an amorphous core on the Ge and GeSn surfaces.
Fig. 2. (a ,b, c) PV-TEM images from sample B1 and (d) FFT pattern of the high resolution TEM image (c), revealing the presence of amorphous material inside the GeSn islands and a 45° rotation between their inner square region and their external sides. (e) Schematic representation of an inverted pyramidal defect deduced from AFM and TEM analysis.
166
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
Fig. 3. Superposition of the line scans taken across representative islands in the b110 N direction for sample A1 (from the AFM image of Fig. 1(a), top panel), showing their evolution from rings to islands.
and Ge concentration profiles (Fig. 4(e)) performed across an inverted pyramidal defect of sample B1 (along the purple line in the HAADFSTEM image of Fig. 4(d)). In addition, Fig. 4(b) shows a local enrichment of Sn atoms (from 15 at.% to 20 at.%) along the sidewalls of the inverted pyramid (represented by the dashed line). Accordingly, Sn-enriched sidewalls are also evident from the 2-D map of the Sn concentration measured in a plane perpendicular to the z direction across the pyramid (Fig. 4(c)). This might be due to an out-diffusion of the Sn atoms from the amorphous phase towards the pyramid walls. However, the Sn atoms inside the amorphous regions do not show diffusing trails under APT inspection. It is worth to notice that the Pt atoms diffused in the amorphous regions have a larger field evaporation value as compared to Ge and Sn atoms. Therefore, in order to enable field evaporation a local smaller radius is formed at the a-GeSn region, thereby
leading to local magnification effects which cause a decrease in magnification at the c-GeSn side of the interface and an increase in magnification at the a-GeSn side of the interface. As a consequence, we observe a false higher decay rate (nm/decade) of the Sn profile in the a-GeSn, corresponding in turn to a false slower variation of the Sn concentration observed in the a-GeSn as compared to the one observed in c-GeSn. Since the concentration of Sn as measured in the amorphous islands is found to be approximately half of that found in the crystalline regions for all the GeSn samples analyzed, we rather believe that either the Sn incorporation rate decreases or the Ge incorporation rate increases in the amorphous phase as compared to the crystalline one. Fig. 5 presents a SSRM 2-D map (a) and a 1-D SSRM line scan survey (b) from the GeSn sample B2, whose xTEM image is shown in Fig. 5(c). It is observed that the amorphous material inside the inverted pyramidal defects has a three order of magnitude higher resistivity than the surrounding crystalline regions. The presence of these amorphous regions can therefore be deleterious for high performance GeSn-based electronic devices. For this reason, in the following section we investigate the possible origin of these defects. The transition from a crystalline to an amorphous phase during low temperature epitaxial growth, generally indicated as epitaxial breakdown, can be initiated by several factors. For example, according to the model of Eaglesham et al. [14], the breakdown mechanism can be the consequence of a kinetic surface roughening followed by an accumulation of defects. Moreover, similar features were observed during Si homoepitaxy by Platen et al. [15], who reported several twodimensional defects propagating from the substrate/film interface to the lower tip of the cones. This possibility has been considered in our case. We observed that the island density for thin GeSn layers on a virtual Ge substrate is comparable to the expected threading dislocation density (i.e. 107–108 cm−2). In addition, the island density increases with the strain relaxation degree of the GeSn layers [7], while the island sides are always aligned along the [110] directions, i.e. the direction of the dislocation cross-hatch pattern lines (visible, for example, in Fig. 1(b), bottom). These findings suggested the existence of a correlation between the formation of the islands and that of the dislocations. However, no extended defect could be observed below or around the
Fig. 4. (a) 3-D APT reconstruction of a tip obtained from sample C1. For a clear visualization, Sn atoms are not included and only 20% of the Ge atoms are shown. The black box and arrow in (a) indicate the position and the direction of the selection for extracting the 1-D concentration profiles plotted in (b). (c) 2-D map (extracted from (a)) of the measured Sn concentration in a plane perpendicular to the z direction. (e) Sn and Ge concentration profiles across an inverted pyramidal defect in sample B1 (determined by EDS along the purple line in the HAADFSTEM image in (d)), showing a lower Sn content in the amorphous region. The dashed line in (a), (b), (c) and (d) represents the transition between the crystalline and amorphous GeSn regions. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
167
Fig. 5. SSRM (a) 2-D map and (b) 1-D line scan survey of a GeSn island sample B2, showing the higher resistivity of the a-GeSn inverted pyramidal defect. (c) xTEM image of the same sample.
amorphous regions in the TEM images of both the Ge and GeSn samples (Fig. 1(b)). In addition, the islands are also present in the case of fully strained GeSn layers grown on a Ge substrate (Fig. 6(a)) or on the surface of Ge layers (Fig. 1 (a) top), where no cross-hatch pattern lines are present, and they are still aligned along [110] directions. This indicates that the islands are not created by threading dislocations generated at GeSn/Ge or Ge/Si(001) interfaces. The observed [7] alignment of the island sides with the cross-hatch pattern lines (i.e. along [110] directions) does not contradict the hypothesis of dislocations eventually nucleating at the island sides on the GeSn layers due to an increased strain energy density. However, this alignment is fortuitous, with the [110] directions simply representing the energetically most favorable directions (being also the directions of the atomic bounds). In fact, the coherent crystalline islands formed due to Stranski Krastanov growth during Ge growth on Si are also aligned along [110] directions. It should be mentioned that the features reported in this work are different from the islands which are commonly observed during the Stranski– Krastanov or Volmer–Weber epitaxial growth modes. In fact, these features are not strain-induced, because they also appear during the homoepitaxy of Ge layers on a Ge substrate, and they are not fully crystalline. In addition, in our case Ge or GeSn 2-D growth occurs together with the island formation, while in the Stranski–Krastanov or the Volmer–Weber this does not happen and a continuous layer is only obtained after overlapping and merging of different islands. Following an analogous reasoning, the islands observed in this work are also different from those observed during the growth of GeSn on
Ge(001) substrates by Molecular Beam Epitaxy (at 155 °C and 6.7 × 10−9 Pa) by Bratland et al. [6]. In fact, also those islands are fully crystalline and they are induced by the increasing strain levels in the GeSn layers with higher Sn concentrations. The crystalline regions of the GeSn islands reported in this work are probably slightly elastically relaxed, but they are not expected to participate significantly to GeSn strain relaxation. Even if no extended defects are observed around or below the inverted pyramidal defects, the TEM analysis indicates that their inner crystalline facets are rough and have relatively steep angles (35° or 71°, as observed for Ge and GeSn, respectively and shown in Fig. 1b). These two conditions might contribute to the observed epitaxial breakdown occurring after a critical height of the previously formed crystalline pit has been overcome. Alternatively, the epitaxial breakdown can be caused by a supersaturation of hydrogen atoms during the growth (as proposed by the model of Thiesen et al. [16]). In this work, H is continuously supplied to the surface during growth, coming from both the adsorbed GeH3 hydrides and from the dilution/bubbling system of the Ge/Sn precursors. As a consequence, a H passivation is expected on the surface of the growing Ge and GeSn layers because of the lower H desorption rate as compared to the Ge and GeSn deposition rates. On the one hand, in fact, a H desorption rate of only 0.26 ML/s is calculated for a Ge surface covered by 1 H monolayer at 320 °C in ultra-high vacuum (following Lee et al. [17]) and an even lower desorption rate can be expected at the growth pressure used in this work (i.e. 101 kPa). On the other hand,
Fig. 6. 3-D PV-AFM images showing the islands as observed for (a) a fully strained GeSn layer grown on a Ge substrate (sample D1) and for two Ge layers grown on a virtual Ge substrate (b) without a pre-EPI bake (sample A2) or (c) using H2 as the carrier gas (sample A3, exhibiting the highest island density).
168
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
we observe that the experimental Ge(Sn) deposition rate at 320 °C and atmospheric pressure ranges between 0.82 ML/s (pure Ge) and 6.60 ML/s (GeSn with 11%), which are significantly higher values than the H desorption rate. In this work, we suggest that the crystalline pits may be caused by the formation of clusters of H and/or contaminant atoms poisoning the available surface sites. In fact, such poisoned areas would be energetically unfavorable for the incorporation of the incoming Ge adatoms, which would then deposit around these areas instead. This phenomenon would result in the formation of pits with crystalline facets over them, as also observed during Si deposition on SiC clusters by Weil et al. [18]. Such hypothesis would be supported by the shape of the ring features in the AFM images of Fig. 1(a) (top), representing the early stages of the island evolution depicted in Fig. 3. When the increasing surface energy of the pit walls makes their further expansion inconvenient, deposition occurs also inside the pits. However, the material deposited on the poisoned areas is expected to be amorphous due to the impossibility to reiterate the crystalline order of the substrate into the growing layer. In order to study if the pits can be formed by surface contamination, SIMS measurements have been done on different GeSn layers grown on virtual Ge substrates. In addition, a TXRF analysis has been done on the surface of a Si wafer exposed to a high SnCl4 flow. All possible contaminants (i.e. Cl, S, F, O, H and C) are found to be below SIMS detection limits and the metal contaminants below TXRF detection limits. However, since the detection limits of these techniques (1 × 1010–3 × 1012 at. cm− 2) are much higher than the density of the islands (7 × 106– 1 × 109 cm−2), we cannot exclude the presence of contaminants inducing the pyramidal defects from these measurements. It is worth to mention that neither APT measurements are adequate to detect such low amounts of contaminants. In fact, APT sensitivity is a function of the total number of atoms in the investigated volume. Since we are looking for only a few contaminants just below the tip of the a-GeSn inverted pyramid, the associated counting statistics are extremely low (~ 22.000 atoms in a volume 103 nm3). The noise level in the mass spectra for this volume is ~ 0.01%. Furthermore, there is a peak overlap between (1) Cl+ and Ge2+, (2) Cl2+ and H2O tail, (3) F+ and H2O tail, (4) S2+ and O+, and (5) S+ and O+ 2 , further reducing the sensitivity to these contaminants. Contaminants (if present) are not expected to come from the starting substrate (e.g. residual C or O), but rather from the growing ambient (e.g. precursor or carrier gases or reactor). In fact, the inverted pyramidal defects have different sizes and nucleate at any depth in the sample (not only at the substrate interface). In addition, the density of the island features on the surface of the GeSn layers with the same nominal Sn content increases with layer thickness. Finally, the absence of a pre-EPI bake has no significant effect on the density of the features, which is found to be 1.1 × 109 cm−2 for sample A1 (Fig. 1(a), top), which did receive a pre-Epi bake before growing a 70 nm thick Ge layer, and 9.2 × 108 cm−2 for sample A2 (Fig. 6(b)), which was grown with exactly the same process conditions, except for the pre-Epi bake which was omitted. It is observed that the feature in sample A2 are somewhat smaller and less concentrated than the features on sample A1. This is tentatively attributed to the existence of an incubation time for sample A2 (which did not receive a pre-Epi back) due to the necessity to desorb the native oxide before epitaxial growth can begin. Such a delay, in fact, would allow less time for the nucleation and growth of the features on sample A2 as compared to the case of sample A1. We notice that the use of H2 as carrier gas instead of N2 increases the island density (from 1.1 × 109 (Fig. 1(a), top) to 3.7 × 109 cm− 2 (Fig. 6(c)). This may indicate that at the low growth temperature used, an accumulation of H atoms on the growing surface acts as nucleation points for amorphous growth, as also suggested in Ref. 15 during Si homoepitaxy. Alternatively, it is also possible that the role of the H atoms is that of lowering the energy barrier for the contaminants diffusion, as determined by Monte Carlo simulations of the surface diffusion
of Ni atoms on Ni(100) [19], for example. It is in fact expected that isolated H or contaminants atoms would not be sufficient to disrupt the epitaxial growth, unless larger clusters are formed. The primary role of H in the formation of the island features would also be in agreement with the observation that the density of these features decreases when the Sn content in the GeSn layers increases (at comparable thicknesses). Since GeSn layers with higher Sn contents are grown using higher SnCl4 partial pressures, their surface is likely more covered by Cl atoms, at the expense of H atoms (a Cl passivation being more favorable than a H one thanks to the higher binding energy with the Ge and Sn atoms on the surface). The Cl atoms could then hinder the surface mobility of the contaminants, as they are found to hinder the surface mobility of Ge atoms in Ref. 20. Alternatively, such a lower H coverage at higher SnCl 4 partial pressures would also explain the lower density of island features if they were simply originating by an accumulation of H atoms. The subsequent lower rate of formation of the island-type features on GeSn as compared to Ge may also explain why smaller features can be observed in the former case at comparable layer thicknesses. In fact, having developed at later stages in the growth, the features on GeSn had less time to grow. Consistently, larger features are observed on thicker GeSn layers grown with longer deposition time (as shown, for example, in Fig. 5(c)). We believe that the origin of these features on the GeSn layers is exactly the same as in the Ge ones, i.e. initiated by the interplay of Ge and contaminants or H atoms. However, the features observed on the GeSn samples generally appear in their “final stage” (referring to Fig. 3), which is with the amorphous phase already nucleated and grown. This is attributed to a faster nucleation of the amorphous material in the case of GeSn as compared to Ge, i.e. to a smaller critical dimension of the rings before they are filled (~1 nm according to our AFM inspections), possibly caused by an increased surface energy when the inverted pyramidal defects have GeSn crystalline inner facets instead of Ge inner facets. The observed 45° rotation between the top region of the inner crystalline facets and the external sides of the islands (Fig. 2) might be explained by the preferential nucleation of crystalline GeSn at the four vertices of the pit created by the interplay of Ge and contaminants or H atoms if its base had a rhombohedral or a squared shape. In fact, these vertices would offer lower energy nucleation sites, as observed during the growth of Ge quantum dots on pre-patterned Si(001) substrate [18,21]. Such a rotation is likely responsible for the observed increase in the angle between the inner facets of the pits and the (001) plane from 35° (for Ge) to 71° (for GeSn), in an attempt to minimize the exposed surface area of the pit walls. The incorporation of higher amounts of Sn atoms on the inner facets of the rotated pits is also expected to be driven by the necessity to reduce their surface energy, thanks to the lower free surface energy of Sn as compared to Ge atoms. This would then explain the Sn enrichment along the inverted pyramidal defects walls observed in both the 1D profile and the 2D map of the Sn concentration measured by APT shown in Fig. 4(b) and (c), respectively. Finally, Fig. 7 shows the PV-AFM images of two Ge samples grown on Ge virtual substrates at 1333 Pa (sample E1, in Fig. 7(a)) and 400 °C (sample F1, in Fig. 7(b)), respectively, while keeping the other growth conditions equal to those used for the reference sample A1. The island features with the amorphous core could not be observed on these two layers. It is then believed that either reducing the total pressure or increasing the temperature during the growth can be effective ways to hinder the formation of the island features. This is possibly due to an associated reduction of the H and/or contaminant atoms on the surface at either lower pressures [22–24] or higher temperatures. Accordingly, no islands were observed by AFM inspection on the surface of two GeSn layers (provided by the Forschungszentrum Julich [25]) grown with Ge2H6 and SnCl4 on the same virtual Ge substrates used in this work, but using both a reduced pressure (6000 Pa) and higher temperatures (380 °C–400 °C).
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
169
Fig. 7. PV-AFM images of sample (a) E1 and (b) F1 Ge samples, showing that reducing the pressure or increasing the temperature can be effective ways to avoid the formation of the islands.
4. Conclusions In summary, we have investigated the characteristics of particular islands-type features with an amorphous core that are developed during the low temperature CVD epitaxial growth of Ge and GeSn layers using Ge2H6. Although further investigations are needed to unambiguously identify the origin of these features, we suggest that they may originate from the formation of clusters of H and/or contaminant atoms during growth. These may initially cause the formation of pits with crystalline rough facets over them, resulting in the ring-shaped islands observed under AFM inspection. Then, when an excess surface energy is overcome, an amorphous phase may nucleate inside these pits and fill them. In order to unambiguously confirm this hypothesis, H/contaminant detection techniques with a sensitivity as low as the density of the islands (106–109 cm− 2) should be used. Reducing the pressure and/or increasing the growth temperature have been proposed as effective ways to suppress the formation of these features, likely due to a reduced presence of H and/or contaminant clusters on the surface. Acknowledgments The authors would like to thank Voltaix and Dow for providing Ge2H6 and SnCl4, respectively. We also thank the IMEC core partners within the imec's Industrial Affiliation Programs on logic, memory, and optical devices. Finally, parts of this work frame within a collaboration funded by the fund for Scientific Research-Flanders and the Japan Society for the promotion of Science with project number VS.018.10N. References [1] R. Pillarisetty, Academic and industry research progress in germanium nanodevices, Nature 479 (2011) 324. [2] J.D. Sau, M.L. Cohen, Possibility of increased mobility in Ge–Sn alloy system, Phys. Rev. B 75 (2007) (045208–1). [3] A. Gassenq, F. Gencarelli, J. Van Campenhout, Y. Shimura, R. Loo, G. Narcy, B. Vincent, G. Roelkens, GeSn/Ge heterostructure short-wave infrared photodetectors on silicon, Opt. Express 20 (2012) 27297. [4] J. Liu, R. Camacho-Aguilera, J.T. Bessette, X. Sun, X. Wang, Y. Cai, L.C. Kimerling, J. Michel, Buffer layers for tensile-strained Ge layers, Ge-on-Si optoelectronics, Thin Solid Films 520 (2012) 3354. [5] Y. Shimura, N. Tsutsui, O. Nakatsuka, A. Sakai, S. Zaima, Control of Sn precipitation and strain relaxation in compositionally step-graded Ge1 − xSnx, Jpn. J. Appl. Phys. 48 (2009) 04C130. [6] K.A. Bratland, Y.L. Foo, T. Spila, H.-S. Seo, R.T. Haasch, P. Desjardins, J.E. Greene, Snmediated Ge/Ge(001) growth by low-temperature molecular-beam epitaxy: surface smoothening and enhanced epitaxial thickness, J. Appl. Phys. 97 (2005) 044904.
[7] F. Gencarelli, B. Vincent, J. Demeulemeester, A. Vantomme, A. Moussa, A. Kumar, H. Bender, J. Meersschaut, W. Vandervorst, R. Loo, M. Caymax, K. Temst, M. Heyns, Crystalline properties and strain relaxation mechanism of CVD grown GeSn, ECS J. Solid State Sci. Technol. 2 (2013) 134. [8] S. Takeuchi, A. Sakai, K. Yamamoto, O. Nakatsuka, M. Ogawa, S. Zaima, Growth and structure evaluation of strain-relaxed Ge1 − xSnx buffer layers grown on various types of substrates, Semicond. Sci. Technol. 22 (2007) S231. [9] F. Gencarelli, B. Vincent, L. Souriau, O. Richard, W. Vandervorst, R. Loo, M. Caymax, M. Heyns, Low-temperature Ge and GeSn chemical vapor deposition using Ge2H6, Thin Solid Films 520 (2012) 3211. [10] B. Vincent, F. Gencarelli, H. Bender, C. Merckling, B. Douhard, D.H. Petersen, O. Hansen, H.H. Henrichsen, J. Meersschaut, W. Vandervorst, M. Heyns, R. Loo, M. Caymax, Undoped and in-situ B doped GeSn epitaxial growth on Ge by atmospheric pressure-chemical vapor deposition, Appl. Phys. Lett. 99 (2011) 152103. [11] G. Wang, F.E. Leys, L. Souriau, R. Loo, M. Caymax, D.P. Brunco, J. Geypen, H. Bender, M. Meuris, W. Vandervorst, M. Heyns, Selective epitaxial growth of germanium on Si wafers with shallow trench isolation: an approach for Ge virtual substrates, ECS Trans. 16 (2008) 829. [12] B. Gault, D. Haley, F. de Geuser, M.P. Moody, E.A. Marquis, D.J. Larson, B.P. Geiser, Ultramicroscopy 111 (2011) 448. [13] A.K. Kambham, A. Kumar, M. Gilbert, W. Vandervorst, 3D site specific sample preparation and analysis of 3D devices (FinFETs) by atom probe tomography, Ultramicroscopy 132 (2013) 65. [14] D.J. Eaglesham, H.-J. Gossmann, M. Cerullo, Limiting thickness hepi for epitaxial growth and room-temperature Si growth on Si(100), Phys. Rev. Lett. 65 (1990) 1227. [15] J. Platen, B. Selle, I. Sieber, S. Brehme, U. Zeimer, W. Fuhs, Low-temperature epitaxial growth of Si by electron cyclotron resonance chemical vapor deposition, Thin Solid Films 381 (2001) 22. [16] J. Thiesen, H. Branz, R.S. Crandall, Explanation of the limiting thickness observed in low-temperature silicon epitaxy, Appl. Phys. Lett. 77 (2000) 3589. [17] J.Y. Lee, J.Y. Maeng, A. Kim, Y.E. Cho, S. Kim, Kinetics of H2 (D2) desorption from a Ge (100)-2 × 1 H (D) surface studied using scanning tunneling microscopy and temperature programmed desorption, J. Chem. Phys. 118 (2003) 22. [18] J.D. Weil, X. Deng, M. Krishnamurthy, Preferential nucleation of Ge islands at selforganized pits formed during the growth of thin Si buffer layers on Si(110), J. Appl. Phys. 83 (1998) 212. [19] K. Haug, N-K. N. Do, Kinetic Monte Carlo study of the effect of hydrogen on the two-dimensional epitaxial growth of Ni(100), Phys. Rev. B 60 (1999) 11095. [20] T.I. Kamins, G.A.D. Briggs, R. Stanley Williams, Influence of HCl on the chemical vapor deposition and etching of Ge islands on Si(001), Appl. Phys. Lett. 73 (1998) 1862. [21] H.-M. Chen, C.-H. Kuan, Y.-W. Suen, G.-L. Luo, Y.-P. Lai, F.-M. Wang, S.-T. Chen, Thermally induced morphology evolution of pit-patterned Si substrate and its effect on nucleation properties of Ge dots, Nanotechnology 23 (2012) 015303. [22] A. Hikavyy, E. Rosseel, S. Dhayalan, L. Witters, H. Mertens, H. Bender, P. Favia, R. Loo, High Ge content SiGe thin films: growth, properties and integration, ECS Trans. 64 (2014) 831. [23] R.P. Donovan, Contamination-Free Manufacturing for Semiconductors and Other Precision Products, CRC Press, 2001. [24] J. Safarian, T.A. Engh, Vacuum evaporation of pure metals, Metall. Mater. Trans. A 44 (2013) 747. [25] S. Wirths, D. Buca, A.T. Tiedemann, B. Holländer, P. Bernardy, T. Stoica, D. Grützmacher, S. Mantl, Epitaxial growth of Ge1 − xSnx by reduced pressure CVD using SnCl4 and Ge2H6, ECS Trans. 50 (2012) 885.
Contents lists available at ScienceDirect
Thin Solid Films journal homepage: www.elsevier.com/locate/tsf
Amorphous inclusions during Ge and GeSn epitaxial growth via chemical vapor deposition F. Gencarelli a,b,⁎, Y. Shimura a,c,d, A. Kumar a,c, B. Vincent a, A. Moussa a, D. Vanhaeren a, O. Richard a, H. Bender a, W. Vandervorst a,c, M. Caymax a, R. Loo a, M. Heyns a,b a
imec, Kapeldreef 75, 3001 Leuven, Belgium Dept. of Metallurgy and Materials Engineering, KU Leuven, B-3001 Leuven, Belgium Nuclear and Radiation Physics Section, KU Leuven, B-3001 Leuven, Belgium d FWO Pegasus Marie Curie Fellow b c
a r t i c l e
i n f o
Article history: Received 10 September 2014 Received in revised form 30 July 2015 Accepted 30 July 2015 Available online 1 August 2015 Keywords: Germanium tin Amorphous Islands Epitaxial breakdown Low temperature chemical vapor deposition Digermane
a b s t r a c t In this work, we discuss the characteristics of particular island-type features with an amorphous core that are developed during the low temperature epitaxial growth of Ge and GeSn layers by means of chemical vapor deposition with Ge2H6. Although further investigations are needed to unambiguously identify the origin of these features, we suggest that they are originated by the formation of clusters of H and/or contaminants atoms during growth. These would initially cause the formation of pits with crystalline rough facets over them, resulting in ring-shaped islands. Then, when an excess surface energy is overcome, an amorphous phase would nucleate inside the pits and fill them. Reducing the pressure and/or increasing the growth temperature can be effective ways to prevent the formation of these features, likely due to a reduction of the surface passivation from H and/or contaminant atoms. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Ge and GeSn alloys are interesting group IV semiconductors with potentially superior transport [1,2] and optical [3,4] properties as compared to conventional Si. Low temperature growth processes are often required in order to fabricate Ge or GeSn heterostructures without undesired strain relaxations, elemental intermixing, dopant diffusion or Sn precipitation [5]. However, low-temperature growth is known to have a series of possible undesired consequences, such as the roughening of the surface [6], the formation of defects or the inclusion of localized or extended amorphous regions [6–8] (often referred to as epitaxial breakdown). Using low temperature chemical vapor deposition (CVD) we demonstrated the epitaxial growth of good quality, dislocation-free thin GeSn layers [9,10] using a high-order, more reactive Ge precursor as compared to standard GeH4 (namely Ge2H6) to achieve sufficiently high growth rates. However, we also observed that when the thickness of these layers increases, their surface is increasingly covered by islands, which we initially attributed to strain relaxation [7]. In this work, we investigate in more details the characteristics of these islands and we discuss the presence of an amorphous core inside them. Since the ⁎ Corresponding author at: imec, Kapeldreef 75, 3001 Leuven, Belgium. E-mail address: [email protected] (F. Gencarelli).
http://dx.doi.org/10.1016/j.tsf.2015.07.076 0040-6090/© 2015 Elsevier B.V. All rights reserved.
presence of these particular features is expected to complicate the fabrication of optical or electrical devices containing Ge or GeSn heterostructures and to have a negative impact on their performances, their formation mechanism is investigated in order to understand and possibly prevent their appearance. 2. Experimental details The analyzed Ge and GeSn layers are epitaxially grown on Ge(001) or on 1 μm Ge-buffered [11] Si(001) substrates in a 200 mm ASMEpsilonTM-like CVD reactor, employing a growth temperature of 320 °C and a growth pressure of 101 kPa (unless otherwise specified). Ge2H6 (diluted in H2) and SnCl4 are used as the Ge and Sn precursors, respectively, with either N2 or H2 as the carrier gas [9]. In order to remove the native oxides, a pre-Epi cleaning of the wafers is performed, consisting of a HF (2%) dip for 90 s followed by a rinsing step in deionized water and a Marangoni Dry step. In addition, the wafers received a standard in-situ pre-Epi bake under H2 for 10 min, at 650 °C and 5333 Pa. The surface morphology of the grown layers is inspected by atomic force microscopy (AFM) measurements, which are executed in tapping mode on an Icon PT tool coupled with a Nanoscope V controller (Bruker), using an AC160TS Olympus probe. The layer crystallinity is characterized using plan-view (PV) and cross-section (X) transmission
164
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
electron microscopy (TEM) together with a fast Fourier transform (FFT) analysis. Atom probe tomography (APT) and energy dispersive X-ray spectroscopy (EDS) are employed to analyze the distribution of the Sn atoms inside the GeSn islands. The APT analysis was done with the Laser Assisted Wide Angle Tomographic Atom Probe (LAWATAP) tool from Cameca using laser pulsing (wavelength 515 nm, 400 fs pulse duration) at 15 K sample temperature. The raw data were reconstructed using TAP3D data software from Cameca using the ‘Voltage curve’ reconstruction method [12]. Tips suitable for APT analysis were prepared by the lift out method and sub-sequential annular Focused Ion Beam (FIB) milling, on a FEI NOVA-600 dual beam tool. The TEM, EDS and high angle annular dark field scanning transmission electron microscopy (HAADF-STEM) analyses are performed on a Tecnai F30 ST tool from FEI using a FEG electron source operated at 300 kV. The X-TEM specimens are prepared by FIB milling (Strata, FEI). A CVD glass, a sputtered Al and a sputtered Pt layers are deposited beforehand on the area of interest. The PV-TEM specimens are prepared using a precision ion polishing system (PIPS, Gatan). The spreading resistance of the GeSn layers is measured in high vacuum (1.3 × 10−3 Pa) by scanning spreading resistance microscopy (SSRM) using a Bruker E-scope with a 50 mV sample bias and a deflection setpoint of 0.5 V. Finally, secondary ion mass spectrometry (SIMS) and total reflection X-ray fluorescence (TXRF) are used to address the incorporation of contaminant atoms inside the GeSn layers. SIMS measurements were performed using an Atomika 4500 tool, with a Cs + primary beam having an incidence angle of 45° and an impact energy of 5 keV. TXRF measurements are performed with an Atomika-FEI TXRF8300 25031 tool equipped with a W X-ray tube, using 50 kV and 55 mA as the X-ray tube voltage and current, respectively, and a measurement angle of 2.3 mrad. It is worth mentioning that the results of all these analysis are consistent when performed for samples of different thickness (in the range 30–1000 nm) and composition (in the range 0–12 at.% Sn). The pictures discussed in this work are purely selected according to their quality in terms of contrast and brightness, for the sake of clarity. The structural details of the corresponding samples are summarized in Table 1.
3. Results and discussion As previously reported [7], particular island-type features are formed on the surface of the growing Ge or GeSn layers during low temperature CVD experiments with Ge2H6. These islands, shown in the top view 2D AFM images of Fig. 1(a) for Ge (sample A1, top panel) and GeSn (sample B1, bottom panel), respectively, have different features, shapes and sizes, as it will be discussed subsequently. In the case of pure Ge growth, the X-TEM inspection (Fig. 1(b), top panel) reveals that inside each crystalline island there is an inverted pyramidal defect containing amorphous material. These defects have their symmetry axis along the growth direction and appear to be delimited by {112} facets, which form an angle of 35° with the (001) growing surface. As it is shown in Fig. 1(b) (bottom
panel), the GeSn islands contain similar defects, but apparently delimited by {221} facets forming an angle of 71° with the (001) plane. Due to the poor contrast between the crystalline and amorphous regions of the GeSn sample B1 in the X-TEM image in Fig. 1 (b) (bottom panel), PV-TEM images with increasing magnification factors are also shown for this sample in Fig. 2(a), (b) and (c). We observe that the top central region of each island has always a square shape, with rounded corners and sides along [010] directions, i.e. 45° rotated as compared to the external island sides aligned along [110] directions. This suggests that the inverted pyramidal defects inside the GeSn islands are actually delimited by {041} crystalline facets, as it is schematically shown in Fig. 2(e). Such a 45° rotation is not observed in the AFM image of the Ge layers (upper panel of Fig. 1(a)). We also observe that the stressinduced bend contours patterns present inside each GeSn island do not continue in its central region (Fig. 2(a) or (b)) and that diffuse rings appear in the FFT image (Fig. 2(d)) obtained from the highest resolution PV-TEM picture (in Fig. 2(c)). These findings indicate the presence of amorphous material inside the crystalline inverted pyramidal defects, also clearly visible in Fig. 2(c). Fig. 3 shows line scans across different islands seen in the top panel of Fig. 1(a) (corresponding to a 70 nm thick Ge layer), but plotted on the same x-axis. By overlapping the line scan profiles in this way, it is possible to observe the evolution of the islands for increasing deposition times. At their beginning stage, they appear as elevated “squared rings” with a dip in the middle (i.e. a pit). Subsequently, at some critical thickness (~ 5 nm), the excess surface energy associated with the pit walls is likely too large and the pit gradually fills in. However, the newly nucleated phase is amorphous and has a higher growth rate than the surrounding crystalline material in the ring, which continues its vertical and lateral growth. APT analysis is also used in order to characterize the amorphous inverted pyramids, which can be localized in the APT tip via the method detailed in Ref. 13. The 3-D reconstruction of a tip obtained from the GeSn sample C1 is shown in Fig. 4(a). For a clear visualization, Sn atoms are not visualized and only 20% of the Ge atoms are shown. We notice that during the final cleaning of the tip at low energies some Pt atoms (coming from a cap previously deposited on the GeSn layer in order to protect it from major ion beam damages) preferentially diffuse towards the amorphous region. This artifact is beneficial to easily visualize the inverted pyramidal defect, whose facets are found to form 71° with the (001) plane, in good agreement with X-TEM results. A 1-D composition profile – extracted from a region across the pyramid walls with the depth aligned along the cross-section of the tip (indicated by a black box in Fig. 4(a)) – is shown in Fig. 4(b). The Sn concentration inside the amorphous pyramid is found to be lower (8–10 at.%) as compared to that in the surrounding crystalline matrix (15 at.%). This variation in the Sn concentration is reflected by a complementary variation in the Ge concentration. It should be mentioned that these APT measurements were performed at a lower pulse fraction than the optimal one necessary to accurately quantify the absolute composition of the GeSn layers [13], which is then slightly overestimated. A lower Sn content in the amorphous GeSn region as compared to the crystalline area is also determined by EDS measurements of the Sn
Table 1 Samples discussed in this work. Sample ID
Epilayer thickness [nm]
Sn content [%]
Carrier gas
Substrate
Comment
A1 A2 A3 B1 B2 C1 D1 E1 F1
70 70 70 144 542 1000 27 295 67
0 0 0 6.7 6.1 12.7 6.4 0 0
N2 N2 H2 N2 N2 N2 N2 N2 N2
1 μm Ge/Si(001) 1 μm Ge/Si(001) 1 μm Ge/Si(001) 1 μm Ge/Si(001) 350 nm Ge/Si(001) 1 μm Ge/Si(001) Ge(001) 1 μm Ge/Si(001) 1 μm Ge/Si(001)
Reference No pre-EPI bake H2 carrier – – – Fully strained 10 Torr 400 °C
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
165
Fig. 1. (a) Top view 2D AFM and (b) X-TEM images of samples A1 (top) and B1 (bottom), indicating the presence of particular island-like features with an amorphous core on the Ge and GeSn surfaces.
Fig. 2. (a ,b, c) PV-TEM images from sample B1 and (d) FFT pattern of the high resolution TEM image (c), revealing the presence of amorphous material inside the GeSn islands and a 45° rotation between their inner square region and their external sides. (e) Schematic representation of an inverted pyramidal defect deduced from AFM and TEM analysis.
166
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
Fig. 3. Superposition of the line scans taken across representative islands in the b110 N direction for sample A1 (from the AFM image of Fig. 1(a), top panel), showing their evolution from rings to islands.
and Ge concentration profiles (Fig. 4(e)) performed across an inverted pyramidal defect of sample B1 (along the purple line in the HAADFSTEM image of Fig. 4(d)). In addition, Fig. 4(b) shows a local enrichment of Sn atoms (from 15 at.% to 20 at.%) along the sidewalls of the inverted pyramid (represented by the dashed line). Accordingly, Sn-enriched sidewalls are also evident from the 2-D map of the Sn concentration measured in a plane perpendicular to the z direction across the pyramid (Fig. 4(c)). This might be due to an out-diffusion of the Sn atoms from the amorphous phase towards the pyramid walls. However, the Sn atoms inside the amorphous regions do not show diffusing trails under APT inspection. It is worth to notice that the Pt atoms diffused in the amorphous regions have a larger field evaporation value as compared to Ge and Sn atoms. Therefore, in order to enable field evaporation a local smaller radius is formed at the a-GeSn region, thereby
leading to local magnification effects which cause a decrease in magnification at the c-GeSn side of the interface and an increase in magnification at the a-GeSn side of the interface. As a consequence, we observe a false higher decay rate (nm/decade) of the Sn profile in the a-GeSn, corresponding in turn to a false slower variation of the Sn concentration observed in the a-GeSn as compared to the one observed in c-GeSn. Since the concentration of Sn as measured in the amorphous islands is found to be approximately half of that found in the crystalline regions for all the GeSn samples analyzed, we rather believe that either the Sn incorporation rate decreases or the Ge incorporation rate increases in the amorphous phase as compared to the crystalline one. Fig. 5 presents a SSRM 2-D map (a) and a 1-D SSRM line scan survey (b) from the GeSn sample B2, whose xTEM image is shown in Fig. 5(c). It is observed that the amorphous material inside the inverted pyramidal defects has a three order of magnitude higher resistivity than the surrounding crystalline regions. The presence of these amorphous regions can therefore be deleterious for high performance GeSn-based electronic devices. For this reason, in the following section we investigate the possible origin of these defects. The transition from a crystalline to an amorphous phase during low temperature epitaxial growth, generally indicated as epitaxial breakdown, can be initiated by several factors. For example, according to the model of Eaglesham et al. [14], the breakdown mechanism can be the consequence of a kinetic surface roughening followed by an accumulation of defects. Moreover, similar features were observed during Si homoepitaxy by Platen et al. [15], who reported several twodimensional defects propagating from the substrate/film interface to the lower tip of the cones. This possibility has been considered in our case. We observed that the island density for thin GeSn layers on a virtual Ge substrate is comparable to the expected threading dislocation density (i.e. 107–108 cm−2). In addition, the island density increases with the strain relaxation degree of the GeSn layers [7], while the island sides are always aligned along the [110] directions, i.e. the direction of the dislocation cross-hatch pattern lines (visible, for example, in Fig. 1(b), bottom). These findings suggested the existence of a correlation between the formation of the islands and that of the dislocations. However, no extended defect could be observed below or around the
Fig. 4. (a) 3-D APT reconstruction of a tip obtained from sample C1. For a clear visualization, Sn atoms are not included and only 20% of the Ge atoms are shown. The black box and arrow in (a) indicate the position and the direction of the selection for extracting the 1-D concentration profiles plotted in (b). (c) 2-D map (extracted from (a)) of the measured Sn concentration in a plane perpendicular to the z direction. (e) Sn and Ge concentration profiles across an inverted pyramidal defect in sample B1 (determined by EDS along the purple line in the HAADFSTEM image in (d)), showing a lower Sn content in the amorphous region. The dashed line in (a), (b), (c) and (d) represents the transition between the crystalline and amorphous GeSn regions. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
167
Fig. 5. SSRM (a) 2-D map and (b) 1-D line scan survey of a GeSn island sample B2, showing the higher resistivity of the a-GeSn inverted pyramidal defect. (c) xTEM image of the same sample.
amorphous regions in the TEM images of both the Ge and GeSn samples (Fig. 1(b)). In addition, the islands are also present in the case of fully strained GeSn layers grown on a Ge substrate (Fig. 6(a)) or on the surface of Ge layers (Fig. 1 (a) top), where no cross-hatch pattern lines are present, and they are still aligned along [110] directions. This indicates that the islands are not created by threading dislocations generated at GeSn/Ge or Ge/Si(001) interfaces. The observed [7] alignment of the island sides with the cross-hatch pattern lines (i.e. along [110] directions) does not contradict the hypothesis of dislocations eventually nucleating at the island sides on the GeSn layers due to an increased strain energy density. However, this alignment is fortuitous, with the [110] directions simply representing the energetically most favorable directions (being also the directions of the atomic bounds). In fact, the coherent crystalline islands formed due to Stranski Krastanov growth during Ge growth on Si are also aligned along [110] directions. It should be mentioned that the features reported in this work are different from the islands which are commonly observed during the Stranski– Krastanov or Volmer–Weber epitaxial growth modes. In fact, these features are not strain-induced, because they also appear during the homoepitaxy of Ge layers on a Ge substrate, and they are not fully crystalline. In addition, in our case Ge or GeSn 2-D growth occurs together with the island formation, while in the Stranski–Krastanov or the Volmer–Weber this does not happen and a continuous layer is only obtained after overlapping and merging of different islands. Following an analogous reasoning, the islands observed in this work are also different from those observed during the growth of GeSn on
Ge(001) substrates by Molecular Beam Epitaxy (at 155 °C and 6.7 × 10−9 Pa) by Bratland et al. [6]. In fact, also those islands are fully crystalline and they are induced by the increasing strain levels in the GeSn layers with higher Sn concentrations. The crystalline regions of the GeSn islands reported in this work are probably slightly elastically relaxed, but they are not expected to participate significantly to GeSn strain relaxation. Even if no extended defects are observed around or below the inverted pyramidal defects, the TEM analysis indicates that their inner crystalline facets are rough and have relatively steep angles (35° or 71°, as observed for Ge and GeSn, respectively and shown in Fig. 1b). These two conditions might contribute to the observed epitaxial breakdown occurring after a critical height of the previously formed crystalline pit has been overcome. Alternatively, the epitaxial breakdown can be caused by a supersaturation of hydrogen atoms during the growth (as proposed by the model of Thiesen et al. [16]). In this work, H is continuously supplied to the surface during growth, coming from both the adsorbed GeH3 hydrides and from the dilution/bubbling system of the Ge/Sn precursors. As a consequence, a H passivation is expected on the surface of the growing Ge and GeSn layers because of the lower H desorption rate as compared to the Ge and GeSn deposition rates. On the one hand, in fact, a H desorption rate of only 0.26 ML/s is calculated for a Ge surface covered by 1 H monolayer at 320 °C in ultra-high vacuum (following Lee et al. [17]) and an even lower desorption rate can be expected at the growth pressure used in this work (i.e. 101 kPa). On the other hand,
Fig. 6. 3-D PV-AFM images showing the islands as observed for (a) a fully strained GeSn layer grown on a Ge substrate (sample D1) and for two Ge layers grown on a virtual Ge substrate (b) without a pre-EPI bake (sample A2) or (c) using H2 as the carrier gas (sample A3, exhibiting the highest island density).
168
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
we observe that the experimental Ge(Sn) deposition rate at 320 °C and atmospheric pressure ranges between 0.82 ML/s (pure Ge) and 6.60 ML/s (GeSn with 11%), which are significantly higher values than the H desorption rate. In this work, we suggest that the crystalline pits may be caused by the formation of clusters of H and/or contaminant atoms poisoning the available surface sites. In fact, such poisoned areas would be energetically unfavorable for the incorporation of the incoming Ge adatoms, which would then deposit around these areas instead. This phenomenon would result in the formation of pits with crystalline facets over them, as also observed during Si deposition on SiC clusters by Weil et al. [18]. Such hypothesis would be supported by the shape of the ring features in the AFM images of Fig. 1(a) (top), representing the early stages of the island evolution depicted in Fig. 3. When the increasing surface energy of the pit walls makes their further expansion inconvenient, deposition occurs also inside the pits. However, the material deposited on the poisoned areas is expected to be amorphous due to the impossibility to reiterate the crystalline order of the substrate into the growing layer. In order to study if the pits can be formed by surface contamination, SIMS measurements have been done on different GeSn layers grown on virtual Ge substrates. In addition, a TXRF analysis has been done on the surface of a Si wafer exposed to a high SnCl4 flow. All possible contaminants (i.e. Cl, S, F, O, H and C) are found to be below SIMS detection limits and the metal contaminants below TXRF detection limits. However, since the detection limits of these techniques (1 × 1010–3 × 1012 at. cm− 2) are much higher than the density of the islands (7 × 106– 1 × 109 cm−2), we cannot exclude the presence of contaminants inducing the pyramidal defects from these measurements. It is worth to mention that neither APT measurements are adequate to detect such low amounts of contaminants. In fact, APT sensitivity is a function of the total number of atoms in the investigated volume. Since we are looking for only a few contaminants just below the tip of the a-GeSn inverted pyramid, the associated counting statistics are extremely low (~ 22.000 atoms in a volume 103 nm3). The noise level in the mass spectra for this volume is ~ 0.01%. Furthermore, there is a peak overlap between (1) Cl+ and Ge2+, (2) Cl2+ and H2O tail, (3) F+ and H2O tail, (4) S2+ and O+, and (5) S+ and O+ 2 , further reducing the sensitivity to these contaminants. Contaminants (if present) are not expected to come from the starting substrate (e.g. residual C or O), but rather from the growing ambient (e.g. precursor or carrier gases or reactor). In fact, the inverted pyramidal defects have different sizes and nucleate at any depth in the sample (not only at the substrate interface). In addition, the density of the island features on the surface of the GeSn layers with the same nominal Sn content increases with layer thickness. Finally, the absence of a pre-EPI bake has no significant effect on the density of the features, which is found to be 1.1 × 109 cm−2 for sample A1 (Fig. 1(a), top), which did receive a pre-Epi bake before growing a 70 nm thick Ge layer, and 9.2 × 108 cm−2 for sample A2 (Fig. 6(b)), which was grown with exactly the same process conditions, except for the pre-Epi bake which was omitted. It is observed that the feature in sample A2 are somewhat smaller and less concentrated than the features on sample A1. This is tentatively attributed to the existence of an incubation time for sample A2 (which did not receive a pre-Epi back) due to the necessity to desorb the native oxide before epitaxial growth can begin. Such a delay, in fact, would allow less time for the nucleation and growth of the features on sample A2 as compared to the case of sample A1. We notice that the use of H2 as carrier gas instead of N2 increases the island density (from 1.1 × 109 (Fig. 1(a), top) to 3.7 × 109 cm− 2 (Fig. 6(c)). This may indicate that at the low growth temperature used, an accumulation of H atoms on the growing surface acts as nucleation points for amorphous growth, as also suggested in Ref. 15 during Si homoepitaxy. Alternatively, it is also possible that the role of the H atoms is that of lowering the energy barrier for the contaminants diffusion, as determined by Monte Carlo simulations of the surface diffusion
of Ni atoms on Ni(100) [19], for example. It is in fact expected that isolated H or contaminants atoms would not be sufficient to disrupt the epitaxial growth, unless larger clusters are formed. The primary role of H in the formation of the island features would also be in agreement with the observation that the density of these features decreases when the Sn content in the GeSn layers increases (at comparable thicknesses). Since GeSn layers with higher Sn contents are grown using higher SnCl4 partial pressures, their surface is likely more covered by Cl atoms, at the expense of H atoms (a Cl passivation being more favorable than a H one thanks to the higher binding energy with the Ge and Sn atoms on the surface). The Cl atoms could then hinder the surface mobility of the contaminants, as they are found to hinder the surface mobility of Ge atoms in Ref. 20. Alternatively, such a lower H coverage at higher SnCl 4 partial pressures would also explain the lower density of island features if they were simply originating by an accumulation of H atoms. The subsequent lower rate of formation of the island-type features on GeSn as compared to Ge may also explain why smaller features can be observed in the former case at comparable layer thicknesses. In fact, having developed at later stages in the growth, the features on GeSn had less time to grow. Consistently, larger features are observed on thicker GeSn layers grown with longer deposition time (as shown, for example, in Fig. 5(c)). We believe that the origin of these features on the GeSn layers is exactly the same as in the Ge ones, i.e. initiated by the interplay of Ge and contaminants or H atoms. However, the features observed on the GeSn samples generally appear in their “final stage” (referring to Fig. 3), which is with the amorphous phase already nucleated and grown. This is attributed to a faster nucleation of the amorphous material in the case of GeSn as compared to Ge, i.e. to a smaller critical dimension of the rings before they are filled (~1 nm according to our AFM inspections), possibly caused by an increased surface energy when the inverted pyramidal defects have GeSn crystalline inner facets instead of Ge inner facets. The observed 45° rotation between the top region of the inner crystalline facets and the external sides of the islands (Fig. 2) might be explained by the preferential nucleation of crystalline GeSn at the four vertices of the pit created by the interplay of Ge and contaminants or H atoms if its base had a rhombohedral or a squared shape. In fact, these vertices would offer lower energy nucleation sites, as observed during the growth of Ge quantum dots on pre-patterned Si(001) substrate [18,21]. Such a rotation is likely responsible for the observed increase in the angle between the inner facets of the pits and the (001) plane from 35° (for Ge) to 71° (for GeSn), in an attempt to minimize the exposed surface area of the pit walls. The incorporation of higher amounts of Sn atoms on the inner facets of the rotated pits is also expected to be driven by the necessity to reduce their surface energy, thanks to the lower free surface energy of Sn as compared to Ge atoms. This would then explain the Sn enrichment along the inverted pyramidal defects walls observed in both the 1D profile and the 2D map of the Sn concentration measured by APT shown in Fig. 4(b) and (c), respectively. Finally, Fig. 7 shows the PV-AFM images of two Ge samples grown on Ge virtual substrates at 1333 Pa (sample E1, in Fig. 7(a)) and 400 °C (sample F1, in Fig. 7(b)), respectively, while keeping the other growth conditions equal to those used for the reference sample A1. The island features with the amorphous core could not be observed on these two layers. It is then believed that either reducing the total pressure or increasing the temperature during the growth can be effective ways to hinder the formation of the island features. This is possibly due to an associated reduction of the H and/or contaminant atoms on the surface at either lower pressures [22–24] or higher temperatures. Accordingly, no islands were observed by AFM inspection on the surface of two GeSn layers (provided by the Forschungszentrum Julich [25]) grown with Ge2H6 and SnCl4 on the same virtual Ge substrates used in this work, but using both a reduced pressure (6000 Pa) and higher temperatures (380 °C–400 °C).
F. Gencarelli et al. / Thin Solid Films 590 (2015) 163–169
169
Fig. 7. PV-AFM images of sample (a) E1 and (b) F1 Ge samples, showing that reducing the pressure or increasing the temperature can be effective ways to avoid the formation of the islands.
4. Conclusions In summary, we have investigated the characteristics of particular islands-type features with an amorphous core that are developed during the low temperature CVD epitaxial growth of Ge and GeSn layers using Ge2H6. Although further investigations are needed to unambiguously identify the origin of these features, we suggest that they may originate from the formation of clusters of H and/or contaminant atoms during growth. These may initially cause the formation of pits with crystalline rough facets over them, resulting in the ring-shaped islands observed under AFM inspection. Then, when an excess surface energy is overcome, an amorphous phase may nucleate inside these pits and fill them. In order to unambiguously confirm this hypothesis, H/contaminant detection techniques with a sensitivity as low as the density of the islands (106–109 cm− 2) should be used. Reducing the pressure and/or increasing the growth temperature have been proposed as effective ways to suppress the formation of these features, likely due to a reduced presence of H and/or contaminant clusters on the surface. Acknowledgments The authors would like to thank Voltaix and Dow for providing Ge2H6 and SnCl4, respectively. We also thank the IMEC core partners within the imec's Industrial Affiliation Programs on logic, memory, and optical devices. Finally, parts of this work frame within a collaboration funded by the fund for Scientific Research-Flanders and the Japan Society for the promotion of Science with project number VS.018.10N. References [1] R. Pillarisetty, Academic and industry research progress in germanium nanodevices, Nature 479 (2011) 324. [2] J.D. Sau, M.L. Cohen, Possibility of increased mobility in Ge–Sn alloy system, Phys. Rev. B 75 (2007) (045208–1). [3] A. Gassenq, F. Gencarelli, J. Van Campenhout, Y. Shimura, R. Loo, G. Narcy, B. Vincent, G. Roelkens, GeSn/Ge heterostructure short-wave infrared photodetectors on silicon, Opt. Express 20 (2012) 27297. [4] J. Liu, R. Camacho-Aguilera, J.T. Bessette, X. Sun, X. Wang, Y. Cai, L.C. Kimerling, J. Michel, Buffer layers for tensile-strained Ge layers, Ge-on-Si optoelectronics, Thin Solid Films 520 (2012) 3354. [5] Y. Shimura, N. Tsutsui, O. Nakatsuka, A. Sakai, S. Zaima, Control of Sn precipitation and strain relaxation in compositionally step-graded Ge1 − xSnx, Jpn. J. Appl. Phys. 48 (2009) 04C130. [6] K.A. Bratland, Y.L. Foo, T. Spila, H.-S. Seo, R.T. Haasch, P. Desjardins, J.E. Greene, Snmediated Ge/Ge(001) growth by low-temperature molecular-beam epitaxy: surface smoothening and enhanced epitaxial thickness, J. Appl. Phys. 97 (2005) 044904.
[7] F. Gencarelli, B. Vincent, J. Demeulemeester, A. Vantomme, A. Moussa, A. Kumar, H. Bender, J. Meersschaut, W. Vandervorst, R. Loo, M. Caymax, K. Temst, M. Heyns, Crystalline properties and strain relaxation mechanism of CVD grown GeSn, ECS J. Solid State Sci. Technol. 2 (2013) 134. [8] S. Takeuchi, A. Sakai, K. Yamamoto, O. Nakatsuka, M. Ogawa, S. Zaima, Growth and structure evaluation of strain-relaxed Ge1 − xSnx buffer layers grown on various types of substrates, Semicond. Sci. Technol. 22 (2007) S231. [9] F. Gencarelli, B. Vincent, L. Souriau, O. Richard, W. Vandervorst, R. Loo, M. Caymax, M. Heyns, Low-temperature Ge and GeSn chemical vapor deposition using Ge2H6, Thin Solid Films 520 (2012) 3211. [10] B. Vincent, F. Gencarelli, H. Bender, C. Merckling, B. Douhard, D.H. Petersen, O. Hansen, H.H. Henrichsen, J. Meersschaut, W. Vandervorst, M. Heyns, R. Loo, M. Caymax, Undoped and in-situ B doped GeSn epitaxial growth on Ge by atmospheric pressure-chemical vapor deposition, Appl. Phys. Lett. 99 (2011) 152103. [11] G. Wang, F.E. Leys, L. Souriau, R. Loo, M. Caymax, D.P. Brunco, J. Geypen, H. Bender, M. Meuris, W. Vandervorst, M. Heyns, Selective epitaxial growth of germanium on Si wafers with shallow trench isolation: an approach for Ge virtual substrates, ECS Trans. 16 (2008) 829. [12] B. Gault, D. Haley, F. de Geuser, M.P. Moody, E.A. Marquis, D.J. Larson, B.P. Geiser, Ultramicroscopy 111 (2011) 448. [13] A.K. Kambham, A. Kumar, M. Gilbert, W. Vandervorst, 3D site specific sample preparation and analysis of 3D devices (FinFETs) by atom probe tomography, Ultramicroscopy 132 (2013) 65. [14] D.J. Eaglesham, H.-J. Gossmann, M. Cerullo, Limiting thickness hepi for epitaxial growth and room-temperature Si growth on Si(100), Phys. Rev. Lett. 65 (1990) 1227. [15] J. Platen, B. Selle, I. Sieber, S. Brehme, U. Zeimer, W. Fuhs, Low-temperature epitaxial growth of Si by electron cyclotron resonance chemical vapor deposition, Thin Solid Films 381 (2001) 22. [16] J. Thiesen, H. Branz, R.S. Crandall, Explanation of the limiting thickness observed in low-temperature silicon epitaxy, Appl. Phys. Lett. 77 (2000) 3589. [17] J.Y. Lee, J.Y. Maeng, A. Kim, Y.E. Cho, S. Kim, Kinetics of H2 (D2) desorption from a Ge (100)-2 × 1 H (D) surface studied using scanning tunneling microscopy and temperature programmed desorption, J. Chem. Phys. 118 (2003) 22. [18] J.D. Weil, X. Deng, M. Krishnamurthy, Preferential nucleation of Ge islands at selforganized pits formed during the growth of thin Si buffer layers on Si(110), J. Appl. Phys. 83 (1998) 212. [19] K. Haug, N-K. N. Do, Kinetic Monte Carlo study of the effect of hydrogen on the two-dimensional epitaxial growth of Ni(100), Phys. Rev. B 60 (1999) 11095. [20] T.I. Kamins, G.A.D. Briggs, R. Stanley Williams, Influence of HCl on the chemical vapor deposition and etching of Ge islands on Si(001), Appl. Phys. Lett. 73 (1998) 1862. [21] H.-M. Chen, C.-H. Kuan, Y.-W. Suen, G.-L. Luo, Y.-P. Lai, F.-M. Wang, S.-T. Chen, Thermally induced morphology evolution of pit-patterned Si substrate and its effect on nucleation properties of Ge dots, Nanotechnology 23 (2012) 015303. [22] A. Hikavyy, E. Rosseel, S. Dhayalan, L. Witters, H. Mertens, H. Bender, P. Favia, R. Loo, High Ge content SiGe thin films: growth, properties and integration, ECS Trans. 64 (2014) 831. [23] R.P. Donovan, Contamination-Free Manufacturing for Semiconductors and Other Precision Products, CRC Press, 2001. [24] J. Safarian, T.A. Engh, Vacuum evaporation of pure metals, Metall. Mater. Trans. A 44 (2013) 747. [25] S. Wirths, D. Buca, A.T. Tiedemann, B. Holländer, P. Bernardy, T. Stoica, D. Grützmacher, S. Mantl, Epitaxial growth of Ge1 − xSnx by reduced pressure CVD using SnCl4 and Ge2H6, ECS Trans. 50 (2012) 885.