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Ge-rich islands grown on patterned Si substrates by low-energy plasma-enhanced chemical vapour deposition This article has been downloaded from IOPscience. Please scroll down to see the full text article. 2010 Nanotechnology 21 475302 (http://iopscience.iop.org/0957-4484/21/47/475302) Download details: IP Address: 131.175.59.76 The article was downloaded on 02/11/2010 at 08:58 Please note that terms and conditions apply. View the table of contents for this issue, or go to the journal homepage for more Home Search Collections Journals About Contact us My IOPscience
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Ge-rich islands grown on patterned Si substrates by low-energy plasma-enhanced chemical vapour deposition

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Page 1: Ge-rich islands grown on patterned Si substrates by low-energy plasma-enhanced chemical vapour deposition

Ge-rich islands grown on patterned Si substrates by low-energy plasma-enhanced chemical

vapour deposition

This article has been downloaded from IOPscience. Please scroll down to see the full text article.

2010 Nanotechnology 21 475302

(http://iopscience.iop.org/0957-4484/21/47/475302)

Download details:

IP Address: 131.175.59.76

The article was downloaded on 02/11/2010 at 08:58

Please note that terms and conditions apply.

View the table of contents for this issue, or go to the journal homepage for more

Home Search Collections Journals About Contact us My IOPscience

Page 2: Ge-rich islands grown on patterned Si substrates by low-energy plasma-enhanced chemical vapour deposition

IOP PUBLISHING NANOTECHNOLOGY

Nanotechnology 21 (2010) 475302 (6pp) doi:10.1088/0957-4484/21/47/475302

Ge-rich islands grown on patterned Sisubstrates by low-energy plasma-enhancedchemical vapour depositionM Bollani1, D Chrastina2, A Fedorov1, R Sordan2, A Picco3 andE Bonera3

1 CNISM and L-NESS, Dipartimento di Fisica del Politecnico di Milano, Polo Regionale diComo, Via Anzani 42, I-22100 Como, Italy2 L-NESS, Dipartimento di Fisica del Politecnico di Milano, Polo Regionale di Como,Via Anzani 42, I-22100 Como, Italy3 Dipartimento di Scienza dei Materiali, and L-NESS, Universita degli Studi diMilano-Bicocca, via Cozzi 53, I-20125 Milano, Italy

E-mail: [email protected]

Received 28 May 2010, in final form 23 September 2010Published 29 October 2010Online at stacks.iop.org/Nano/21/475302

AbstractSi1−x Gex islands grown on Si patterned substrates have received considerable attention duringthe last decade for potential applications in microelectronics and optoelectronics. In this workwe propose a new methodology to grow Ge-rich islands using a chemical vapour depositiontechnique. Electron-beam lithography is used to pre-pattern Si substrates, creating materialtraps. Epitaxial deposition of thin Ge films by low-energy plasma-enhanced chemical vapourdeposition then leads to the formation of Ge-rich Si1−x Gex islands (x > 0.8) with ahomogeneous size distribution, precisely positioned with respect to the substrate pattern. Theisland morphology was characterized by atomic force microscopy, and the Ge content and strainin the islands was studied by μRaman spectroscopy. This characterization indicates a uniformdistribution of islands with high Ge content and low strain: this suggests that the relatively highgrowth rate (0.1 nm s−1) and low temperature (650 ◦C) used is able to limit Si intermixing,while maintaining a long enough adatom diffusion length to prevent nucleation of islandsoutside pits. This offers the novel possibility of using these Ge-rich islands to induce strain in aSi cap.

(Some figures in this article are in colour only in the electronic version)

1. Introduction

Germanium-rich Si1−x Gex quantum dots have been studiedand analysed in the past few years for microelectronic andoptoelectronic applications [1–6] with the idea of enhancingthe electron or hole mobility in Si substrates by locallyinducing stress. Since Ge islands tend to nucleate randomlyon flat Si substrates, patterning of the substrate can be used toachieve controlled positioning by creating material traps [7].Here, we use electron-beam lithography (EBL) and reactiveion etching (RIE) to realize ordered arrays of pits, with varioussize and periodicity, on Si(001) substrates. The subsequentepitaxial deposition of a thin film of Ge by low-energy

plasma-enhanced chemical vapour deposition (LEPECVD) [8]results in the nucleation of Ge-rich islands on top of the pits.LEPECVD has been used to grow high quality strain-relaxedbuffers [9], strained two-dimensional Ge quantum wells, [10]and Ge-rich multiple quantum well structures [11]; here weemploy LEPECVD to obtain very thin Ge layers in order tocreate three-dimensional structures. However, in addition tothe precise positioning of Si1−x Gex islands, a control of theirchemical composition is required, since the 3-d compositionprofile ultimately determines their structural and electronicbehaviour and optical properties.

The literature presents different intermixing mechanismsregarding the growth and the evolution of Ge islands in the

0957-4484/10/475302+06$30.00 © 2010 IOP Publishing Ltd Printed in the UK & the USA1

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Nanotechnology 21 (2010) 475302 M Bollani et al

Figure 1. A 10 × 10 μm2 atomic force microscopy (AFM) 3-dimage of self-assembled SiGe islands on a flat Si(001) substratefollowing the deposition of 3.0 nm of Ge at 650 ◦C. The growth ratewas 0.1 nm s−1.

presence of a surface with extrinsic morphology [12–14].In general it has been shown that Si1−xGex islands grownon patterned areas have larger volumes than those onthe corresponding planar substrates, suggesting a differentcomposition state with large intermixing and relaxation onthe patterned substrates [7, 14, 15]. Recent works [16, 17]introduce a quantitative analysis of Ge content in the islandson patterned or flat areas obtained by combining selective wetchemical etching with atomic force microscopy (AFM) andanomalous x-ray scattering, reporting a maximum Ge valueof about 40%, increased to about 50% by stacking [6]. Inthe present work we have used an optimized growth rate tominimize intermixing, and thereby obtain Ge-rich islands. TheGe content and deformation state of the islands inside andoutside the patterned regions was evaluated using μRamanspectroscopy. The measured values are in good agreement withthe volumes needed for pit refilling and island formation asobtained by modelling the AFM profiles.

The technological relevance of results presented in thiswork is considerable, regarding the possibility of obtaining Ge-rich islands on periodic substrates by a CVD technique. Themost important possible application is mobility enhancementin nanofabricated electronic devices. Indeed, such structures,with controllable size and position, can be used to createlocally strained areas in the Si substrate or cap layer in whichthe electron or hole mobility would be increased [6, 18]. SiGehas also been studied for thermoelectric applications [19–21],and Ge-rich SiGe islands present interesting new possibilitiesin this field. SiGe islands also demonstrate interesting opticalproperties [22, 23].

2. Experimental procedure

2.1. Substrate preparation and film deposition

Pit-patterned substrates were prepared by EBL on a Si(001)surface cleaned by an ultrasonic treatment with organicsolvents [24]. During EBL, different doses were used tocreate pits with final diameters of 80–205 nm. For everyexposure, the patterned area was 25 × 25 μm2, with pits ata regular spacing of 1 μm. By RIE, roughly 70 nm of Si wereetched. Before growth, all samples were RCA cleaned [25],

Figure 2. The aspect ratios (ARs) of islands grown on flat(unpatterned) Si(001) substrates as a function of the square root ofthe base area A1/2

0 . Ge was deposited at 0.1 nm s−1 at a substratetemperature of 650 ◦C.

and the native oxide was removed by a HF dip and waterrinse. Epitaxy of Ge on Si(001) by LEPECVD, at the particulargrowth rate (0.1 nm s−1) and substrate temperature (650 ◦C)employed here, proceeds in a Stranski–Krastanov fashionas for molecular-beam epitaxy (MBE) or ultrahigh vacuumCVD [26].

2.2. Characterization

The distribution and size of the islands was characterized byAFM, while the Ge content and strain in individual islandswas measured using μRaman spectroscopy. All AFM analyseswere carried out in tapping mode using an Innova-Veecoinstrument with ultra-sharp tip of typical curvature radius2 nm. The μRaman spectra were acquired in a back-scatteringconfiguration with a 532 nm excitation wavelength in order toobtain a high signal-to-noise ratio with Ge-rich structures. Dueto self-absorption, the investigated depth corresponds to aboutone half of the penetration depth dp, ranging from about 10 to40 nm in Ge-rich Si1−xGex alloys. The microscope features a100×/0.90 NA objective with a spatial resolution of 360 nm.

3. Result and discussion

3.1. Growth on unpatterned Si(001)

An initial set of experiments was used to characterize islandformation on flat Si(001) substrates following the depositionof Ge by LEPECVD. AFM images (not shown) of islandsformed after the deposition of Ge at 400–650 ◦C indicate thatthe island size and spacing increase as temperature increases,for equal thicknesses of deposited Ge. At a fixed substratetemperature of 650 ◦C and growth rate of 0.1 nm s−1, 1–3 nmof Ge were deposited directly on flat, unpatterned substrates.Figure 1 presents a typical AFM image of islands formed afterthe deposition of 3 nm of Ge, in which the islands of various

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Nanotechnology 21 (2010) 475302 M Bollani et al

Figure 3. AFM images of pit refilling after depositing (a) 1.0, (b) 2.0, (c) 2.5 and (d) 3.0 nm pure Ge in which the original diameter is∼205 nm and the depth is ∼70 nm. After 1.0 nm there are four visible lobes inside the pit along the parts of the walls which form {110}surfaces. For 2.0–2.5 nm Ge, islands nucleate at the bottom of the pits and, as a consequence, the pits fill. For 3.0 nm Ge the pit is completelyfilled, and an island has formed with an AR of 0.20.

Figure 4. AFM images of the 80 nm pits show that island growth starts once 2.0 nm of pure Ge has been deposited (a). There are no islandsbetween the pits, and not all pits of the matrix are capped with islands. Panel (b) shows a single island with an AR of 0.16.

sizes are distributed randomly. The aspect ratio (AR) is definedas the height h divided by the square root of the base area,A1/2

0 , and is shown as a function of the A1/20 of the islands in

figure 2. After the deposition of 1 nm of Ge, small islands(A1/2

0 ∼ 0.05–0.10 μm) with ARs of 0.10–0.15, typical ofpyramids [27], are found. After 2 nm, the island populationdivides into two, with one population featuring wide, flatislands which may be dislocated (A1/2

0 up to ∼0.3 μm with AR∼0.15) and the other featuring narrow, taller islands (A1/2

0 ∼0.05–0.10 μm with AR up to ∼0.20). After 3 nm of Ge, mostislands tend towards an AR of 0.20, typical of domes [27].

3.2. Growth on pit-patterned substrates

After 1 nm of pure Ge, the pits are not completely filled:AFM images (figure 3) show that Ge is found on the{110} walls of the pits. Other works in the literature showthat nucleation starts from the bottom of these pits anddome-shaped islands are formed with identical facets, as onunpatterned substrates [26, 28–30]. In this work the nucleationis different: we suggest that growth along the pit walls occursbecause there is no Si buffer layer to shape the indentedwalls. After RIE, in fact, the pit wall is not perfectly flat butpresents nanoscopic steps which happen to be optimal for the

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Nanotechnology 21 (2010) 475302 M Bollani et al

Figure 5. AFM images of the 200 nm pits show that for this matrix it is necessary to deposit 3.0 nm of pure Ge to nucleate islands (a).Ordering and homogeneous size distributions are obtained. Panel (b) shows a single island with an AR of 0.19.

nucleation of the first Ge adatoms to be deposited, showinggrowth faceting along {105} crystallographic planes.

From theoretical computations (not shown), it turns out tobe more convenient for Ge to grow along the vertical walls ofthe pit, in comparison with the flat (001) substrate or bottom ofthe pit [31]. The lobes on the {110} pit walls are able to relaxelastically without inducing strain in the substrate. Therefore,for the same Ge volume, the elastic energy density (calculatedas the ratio of total energy and Ge volume of the system as afunction of pit filling) is lower when the pit is filled from thesidewalls and not from the bottom.

In this work we present only cases where the pits trapall the deposited Ge without the formation of small islandsbetween pits, meaning that the Ge adatom migration lengthL = 2

√Dt (with D the diffusion constant) is much longer

than the pit periodicity a. This result has been obtained afteroptimization of the growth conditions (temperature and growthrates in particular) as a function of the patterning geometry.This means that for a given thickness of deposited material theisland volume depends only on the pit periodicity.

Once the pit is completely filled, an island starts to form.For the same Ge thickness, the behaviour of the deposited Gechanges as a function of the pit diameter. As shown in figure 4,the first islands on the pits with 80 nm diameter appear after2.0 nm of Ge, while for 200 nm pits the islands are found after3.0 nm (figure 5). An island is shown in the right-hand panelof figure 5 in which {1 0 5}, {1 1 3}, and {15 3 23} facets canbe seen. These facets are in accordance with those previouslyreported in the literature [27, 32, 33]. We also observe that thetop of the island is parallel to (001). Figure 6 shows the ARas function of V or A1/2

0 of the islands grown on 80 nm pits,as compared to the data in figure 2 for islands grown on flatsubstrates. It can be seen that islands grown on 80 nm pits aremuch more uniform than those grown on flat substrates. TheARs of islands grown on patterned regions are well defined,being 0.14–0.16 after 2.0 nm Ge and 0.18–0.20 after 3.0 nm.Islands on pits form a relatively homogeneous size distribution,similar to the larger islands nucleated on flat Si(001). After2.0 nm Ge, islands on pits are also flatter (lower AR) than mostof the islands on flat Si(001).

The Raman spectra from two selected nanostructures areshown in figure 7. This measurement was performed on apit of 180 nm diameter which was first partially filled (black

Figure 6. Islands on flat (red circles) and on 80 nm pits (blacksquares) following the deposition of (a) 2.0 and (b) 3.0 nm of Ge at650 ◦C. The ARs of islands grown on patterned regions are welldefined, being 0.14–0.16 after 2.0 nm Ge and 0.18–0.20 after 3.0 nm.Islands on pits form a relatively homogeneous size distribution,similar to the larger islands nucleated on flat Si(001). After 2.0 nmGe, islands on pits are also flatter (lower AR) than most of the islandson flat Si(001).

line), and then capped by an island (red line). From thesespectra one can observe the Ge–Ge mode at about 300 cm−1

and the Si–Ge mode at about 400 cm−1. The Si–Si band in

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Nanotechnology 21 (2010) 475302 M Bollani et al

Figure 7. (a) An example Raman spectra of the investigated structures, showing the Ge–Ge mode peak at about 300 cm−1 and the Si–Gemode peak at about 400 cm−1. The strong peak at about 520 cm−1 is the Si–Si mode peak of the substrate. The spectra were collected fromthe structures grown on 180 nm diameter pits. The black line refers to the partially filled pit, after the deposition of 2.0 nm of Ge. The red linerefers to the island grown on top of the pit, after a 3.0 nm growth of Ge. (b) With the same colour code as panel (a), this panel comparesRaman shifts due to in-plane strain and Ge content in order to determine the values of strain and Ge content as described in the text.

the nanostructures is too weak compared to the 520.5 cm−1

Si–Si band from the substrate of unstrained bulk silicon.The focus was positioned accurately on the centre of thenanostructure using a piezoelectric positioner. The spectralposition of the Ge–Ge and Si–Ge vibrational bands was used tomeasure strain and Ge content [34, 35]. For each spectrum theinvestigated volume can vary significantly with the Ge content.For example, different values of the Ge content x = 1.0,x = 0.8, and x = 0.6 yield values of dp/2 of about 7, 10,and 30 nm respectively. Therefore, it is possible to quantifythe investigated volume of the nanostructure only after themeasurement, when the Ge content is known.

The same procedure was repeated for different nanostruc-tures. Table 1 is a summary of their strain and Ge contents.The data reported were acquired from nanostructures close tothe centres of the patterned grids of pits. We also measurednanostructures at the borders of the patterned regions, but wedid not observe significant differences. As one can infer fromtable 1, the Ge content is always very high as compared to is-lands obtained by MBE, in which the Ge content is usuallybelow 0.40 [16]. Inside the pits, after the deposition of 2.0 nmof Ge, the Ge content is around 0.75; the Ge content increasesto more than 0.90 when 3.0 nm of Ge are deposited. This highGe content may be due to the relatively high growth rate of0.1 nm s−1 obtained by LEPECVD, meaning that the Ge de-position takes only 20–30 s and there is therefore relativelylittle opportunity for intermixing at 650 ◦C. All the structurespresent similar x independently of the diameter d of the pit.The level of strain is always very low, around a few parts perthousand. The data reported were obtained under the assump-tion of a biaxial strain, although it has been shown [5] that theRaman modes of a thin layer with the curved shape of an islandare slightly different with respect to flat heteroepitaxial films.The low level of strain must be attributed to the presence ofdislocations due to the high mismatch between the substrateand the Ge-rich nanostructures. It is also interesting to notice

Table 1. Ge content x and in-plane strain ε‖ for nanostructures ofvarious deposition thicknesses h and pit diameters d . The error barsrefer to the single measurement. We have calculated the in-planestress, σ‖, assuming zero out-of-plane stress and a linearinterpolation of the lattice constants between those of Si and Ge [36].Based on the Ge content measured, an estimation of the investigateddepth is given in the column reporting dp/2. The last column reportsif the structure was a partially filled pit or an island.

h(nm)

d(nm) x

ε‖(×10−3)

σ‖(GPa)

dp/2(nm) Structure

2.0 80 0.82 ± 0.02 3 ± 1 0.5 ± 0.2 12 Island2.0 180 0.76 ± 0.03 −1 ± 1 −0.2 ± 0.2 20 Pit2.0 200 0.75 ± 0.03 −1 ± 1 −0.2 ± 0.2 21 Pit

3.0 80 0.92 ± 0.01 4 ± 1 0.6 ± 0.2 9 Island3.0 180 0.94 ± 0.01 4 ± 1 0.6 ± 0.2 8 Island3.0 200 0.90 ± 0.01 3 ± 1 0.4 ± 0.2 9 Island

that on top of the island the investigated layer is slightly butdetectably strained, with a tensile deformation. This could bedue to the thermal expansion coefficient mismatch between theGe-rich nanostructures and the Si substrate [37, 38], meaningthat the islands were fully relaxed at the growth temperatureand became strained on cooling to room temperature. Semi-filled pits, however, present slight compressive strain, possiblyindicating that this material was not fully relaxed at the growthtemperature or was free to relax during cooling.

4. Conclusions

We have reported for the first time the formation of veryGe-rich Si1−x Gex (x > 0.8) islands grown on pre-patternedSi(001) substrates by a CVD technique. Firstly, we haveinvestigated how Ge growth occurs along the pit walls inthe absence of a Si buffer layer, in terms of energeticconsiderations. Then, we have studied the formation,morphology and composition of the islands as a function of

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Nanotechnology 21 (2010) 475302 M Bollani et al

the thickness of Ge deposited and as a function of substrate pitdiameter. Finally, μRaman analyses have confirmed the highGe content and low strain state of individual islands nucleatedon patterned pits. We speculate that the relatively high growthrate (0.1 nm s−1) and low temperature (650 ◦C) used are able tolimit Si intermixing, while maintaining a long enough adatomdiffusion length to prevent nucleation of islands outside pits.

Acknowledgments

The Cariplo Foundation is gratefully acknowledged for par-tially financing this research through the project MANDIS. Theauthors thank Luca Gagliano for the Raman measurements.

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