Article pubs.acs.org/JPCC
Work Function Measurement of Silicon Germanium Heterostructures Combining Kelvin Force Microscopy and X‑ray Photoelectron Emission Microscopy Sylvain Pouch,†,⊥ Michele Amato,‡ Matteo Bertocchi,§ Stefano Ossicini,§ Nicolas Chevalier,†,⊥ Thierry Mélin,∥ Jean-Michel Hartmann,†,⊥ Olivier Renault,†,⊥ Vincent Delaye,†,⊥ Denis Mariolle,†,⊥ and Łukasz Borowik*,†,⊥ †
Université Grenoble Alpes, F-38000 Grenoble, France CEA, LETI, MINATEC Campus, F-38054 Grenoble, France ‡ Institut d’Electronique Fondamentale, CNRS-UMR 8622, Université Paris-Sud, 91405 Orsay, France § Dipartimento di Scienze e Metodi dell’Ingegneria, Università di Modena e Reggio Emilia, Pad. Morselli, Via Amendola 2, 42122 Reggio Emilia, Italy ∥ Institut d’Electronique de Microélectronique et de Nanotechnologie, CNRS-UMR 8520, Avenue Poincaré, BP 60069,59652 Villeneuve d’Ascq Cedex, France ⊥
ABSTRACT: Work function in Si1−xGex heterostructures with Ge content in the 6% to 49% range was studied with high energy resolution by combining Kelvin force microscopy and X-ray photoelectron emission microscopy. Although the two methods are based on distinct physical mechanisms, we show that both techniques give the same work function differences between each Si1−xGex layer, as small as 20 meV. To detect such small work function differences, we put in evidence the necessity of preparing the Si1−xGex sample surface with polishing, HF etching and Ar+ sputtering. Such surface preparation allows, in principle, to reduce the deleterious influence of surface states, coming for instance from carbon atoms or native oxide, on quantitative work function extraction. We show in this paper that even after such a sample surface preparation, a strong band bending can be present, which causes a contrast inversion on the surface of the material and yields an artificially lower surface work function with respect to theoretical values. By using density functional theory simulations, we demonstrate that such inversion is likely due to residual carbon present on the surface.
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INTRODUCTION Silicon−germanium heterostructures are nowadays used in mass-production of microelectronics circuits. The hole mobility in p-type metal oxide semiconductor field effect transistors (MOSFETs) is indeed several times higher in a compressively strained silicon germanium channel than a pure silicon one. The threshold voltage is also tunable by the Ge content and strain in such channels, which is another advantage.1,2 In situ boron-doped embedded or raised sources and drains can uniaxially strain the channel of very short gate length pMOSFETs, giving another hole mobility boost.2,3 An improved understanding of the electronic properties of Si1−xGex-based heterostructures is desirable, especially concerning the work function (WF) which plays a crucial role in device technology. Theoretical study of the WF in silicon−germanium heterostructures has already been performed by Sant et al.,4 using empirical methods and an interpolation scheme, giving access to the bulk band structure of Si1−xGex as a function of the Ge content x. Moreover, Sant et al.4 proposed an accurate model taking into account the influence of strain on band structure, leading to a better understanding of physics in © 2015 American Chemical Society
Si1−xGex-based devices. However, to the best of our knowledge, no spatially resolved WF study was performed on samples with a large Ge composition range (e.g., 6%−49%, as here). We thus present here, an experimental investigation of Si1−xGex electronic structure with two highly sensitive characterization techniques: Kelvin force microscopy (KFM) and X-ray photoelectron emission microscopy (XPEEM). Both are able to map a sample’s WFa fundamental value directly linked to the position of the Fermi levelat the nanoscale and with an energy sensitivity close to the meV range. However, the physical phenomena exploited by these techniques are quite different. This makes such a cross analysis very interesting. Additional data−like roughness for KFM and chemical composition for XPEEMare also recorded to be used as additional information for the interpretation of WF measurements. Received: September 23, 2015 Revised: November 3, 2015 Published: November 4, 2015 26776
DOI: 10.1021/acs.jpcc.5b09278 J. Phys. Chem. C 2015, 119, 26776−26782
Article
The Journal of Physical Chemistry C The paper is organized as follows: first, the Si1−xGex heterostructure studied, its surface preparation as well as the KFM and XPEEM measurement protocols used are presented. Then, density functional theory (DFT) calculations for the evaluation of the WF in Si and Ge films (pure and mixed with different surface terminations) are described together with complementary results obtained with other surface characterization techniques−like X-ray photoelectron spectroscopy (XPS), Auger electron spectroscopy (AES), transmission electron microscopy (TEM) and TEM energy-dispersive Xray spectroscopy (EDX). Results obtained with KFM and XPEEM are then discussed. WF mapping will notably be compared with calculated WF for so-called “clean” systems. Differences observed between theoretical and measured WF will be discussed, and finally, simulations on structures with residual atomic species at the surface will be presented. It will be shown how residual carbon on Si1−xGex surface can greatly modify the WF.
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EXPERIMENTAL SECTION Epitaxial heterostructures consisting of 800 nm thick Si1−xGex layers stacked one upon the other with increasing Ge content from 6% up to 49% were used in this study. An Epi Centura reduced pressure−chemical vapor deposition (RP-CVD) industrial cluster tool from Applied Materials was used to grow the SiGe multilayers at high temperatures on p-type (1015 cm−3) boron-doped [∼7−10 Ω·cm] Si(001) substrates. The fabrication process is described elsewhere.5 We grew two samples with different Ge contents: from 6% to 29% (sample A) and from 34% to 49% (sample B). In the case of sample B, a rather thick (∼1500 nm) linearly graded SiGe layer was inserted between the Si substrate and the first Si0.66Ge0.34 layer of the actual stack. Its purposes were multiple: (i) confine the misfit dislocations necessary to accommodate the lattice parameter mismatch between Si and the high Ge content stack on top in it, (ii) minimize the roughness of the growing surface and achieve more abrupt interfaces between each Si1−xGex layer of the stack of actual interest, and (iii) maximize the degree of strain relaxation, thereby yielding almost strainfree layers on top. As described in Figure 1a, samples A and B are glued together using a carbon adhesive sheet. Such a bonding prevents edge artifacts coming from parasitic photoemission from the side of the sample.6 In the case of scanning probe microscopy, it also avoids any perturbation of the electric field at the tip due to the presence of a cleaved edge and thus changing electrical capacitances.7 After being glued, samples were polished using a Gatan Centar mechanical polishing machine in order to minimize the surface roughness of the cross-section surface and to reduce the height difference between samples A and B. Topography was checked by atomic force microscopy (AFM) using an Omicron Nanotechnology VT-AFM system with a Nanonis controller (SPECS Zürich). As silicon germanium is harder than silicon, after polishing we measure the presence a rather small (∼10 nm) dome over the large (10 μm) area.8 Such a small topography difference is not expected to affect the WF measurements. The Ge content was quantified by X-ray diffraction (XRD) measurements9 and were cross-checked by local Auger spectroscopy (PHI 700Xi Auger nanoprobe). Before Auger analysis, the sample surface was sputtered with Argon for 20 min to remove surface contamination. A beam energy of 20 keV and a filament current of 1 nA were used for Auger analysis which allows to obtain a lateral resolution of 8 nm for chemical
Figure 1. (a) Schematic representation of the Si1−xGex cross-section studied, where dark gray and bright gray bars represent Ge and Si concentrations, respectively. The so-called “sacrificial layer” is a linearly graded layer used to properly accommodate the lattice mismatch between the Si and layer no. 6. (b) Ge and Si concentration with respect to defined layer numbers. Comparison between EDX and AES. (c) XPS measurements on Si1−xGex sample before and after Ar sputtering. (d) TEM image of the Si0.51Ge0.49 layer cross-section after low energy Ar+ sputtering.
mapping. As shown in Figure 1b, the Si and Ge concentrations from Auger spectroscopy are in good agreement with XRD values. A dip during 30 s in a diluted (10%) hydrofluoric (HF) bath was used to remove the native oxides present on the surface of the SixGe1−x layers. The sample was loaded afterward into an ultrahigh vacuum chamber. The short duration of such a transfer, less than 20 min, enabled us to preserve the H passivation of the surface and prevent reoxidation.10,11 X-ray photoelectron spectroscopy (XPS) (not shown here) confirmed that SiO2 and GeOx were indeed efficiently removed. However, a large amount of carbon, which can influence the SixGe1−x WF measurements, was still present on the surface. 26777
DOI: 10.1021/acs.jpcc.5b09278 J. Phys. Chem. C 2015, 119, 26776−26782
Article
The Journal of Physical Chemistry C
XPEEM analyzer.16 Since φanalyzer is stable during measurements and can be calibrated, the measured φsample is then an absolute value. XPEEM measurements were performed with a NanoESCA spectromicroscope (which is described in more details elsewhere17), which is equipped with low (Hg: hν = 4.9 eV) and high energy excitation sources (HeI VUV and Al Kα Xrays). In order to map WF, series of photoelectron images were acquired for photoelectron energies in the 2−7 eV range (e.g., the secondary electron threshold tail). Here we use 0.025 eV energy steps and acquisition times of 300 s per image. Then, the energy dependent intensity curve for each pixel in the image series was fitted using a complementary error function to obtain the local WF value that was used to obtain the WF map.18 The error in the fitting of photoemission threshold yields a typical sensitivity of 25 meV concerning WF determination. We adjust the pass energy to 50 eV, the entrance slit to 1 μm and the contrast aperture to 150 μm. A lateral resolution of 150 nm is then achieved. Experiments are performed with a 24 × 24 μm2 field of view. The Schottky effect (which decreases the measured surface WF by 0.098 eV due to the high electric field of 6.67 kV/mm) was compensated when producing the final WF maps.15 The local WF determination using XPEEM presupposes that the spectrometer WF is known. Consequently, the measured photoemission threshold energy gives an absolute WF value. For this reason, we will use absolute XPEEM values as references for the relative value KFM measurements. This will be discussed more precisely later on. The calibration of the Nano Esca analyzer has been performed using two crystalline metal surfaces thoroughly prepared in ultrahigh vacuum: Cu [111] and Ag [100]. The measurements of photoemission thresholds were compared with theoretical WF values and the difference between these values gives the spectrometer WF. In order to obtain a precise value of the absolute WF, we did again this calibration after the analysis in order to be sure that the analyzer WF did not change overtime.
Ar+ sputtering (500 eV, 30 min) was thus used to get rid of carbon surface contamination. XPS survey spectra were recorded to verify the presence of oxygen and carbon. Their low intensity confirms that carbon contamination had been eliminated and oxygen contamination considerably reduced (Figure 1c). In addition, we did not notice any argon contamination with XPS. Additionally, TEM and TEM EDX analysis were performed on sample cross sections before and after Ar+ sputtering. Although low energy Ar+ sputtering was used, there are notable structural differences after such a treatment. The first 5 nm away from the sample Si1−xGex surface is amorphous irrespectively of the Ge content (look Figure 1d for example on Si0.51Ge0.49). Additionally, we detect constant concentations of oxygen (∼7%) and argon (∼2%) in the volume of this amorphous layer. WF measurements were performed using two different techniques: KFM, which is a local scanning probe method, and XPEEM, which is a spatially resolved photoemission technique. Both techniques are presented hereafter. KFM is one of the techniques used to detect contact potential difference (CPD), which is equal to the WF difference between the sample and the tip. In KFM, an electrical excitation (dc + ac voltage) is applied to the tip at the cantilever resonance frequency. The electrical excitation generates a cantilever oscillation (or resonance frequency shift) if the cantilever dc bias does not match the sample CPD. To measure CPD, a dc bias feedback is introduced in order to nullify the cantilever oscillation (amplitude modulation mode) or frequency shift (frequency modulation mode) at the angular frequency. This generates a CPD map. During the experiments, both amplitude modulation (AM) and frequency modulation (FM) modes are used. The CPD is directly related to the WF CPD. |e| = φtip − φsample (see ref 12 for further information). This means that our values will be relative WF values. One can calibrate the WF of the tip to obtain theoretical values. However, each tip has a different WF which can also change overtime (damages, tip contamination). KFM measurements have been performed with an Omicron Nanotechnology VT-AFM system with a Nanonis controller (SPECS Zürich) and a Budget Sensors Cr/Pt coated EFM tip (stiffness ∼2 N/m). Here, we used AFM and KFM modes simultaneously. In the AFM mode, the tip is excited at its fundamental resonant frequency, e.g. 67 kHz. The vibration amplitude is set at 100 mV and the resonant frequency shift at −25 Hz. Using the method derived from the work of Giessibl,13 we estimate that the peak-to-peak vibration amplitude is close to 40 nm. Both AM KFM and FM KFM modes were called upon. The first harmonic at 421 kHz is used in the AM mode, with an ac voltage electrical excitation of 100 mV. In the FM mode, we tune the resonant frequency with a 1 kHz modulation and a 500 mV ac voltage.12 XPEEM WF measurement is based on different physical phenomena than KFM. Here, an electron in a bound state is excited by a photon of energy hν. This electron is detected in its final state outside the sample at a kinetic energy. WF can be determined using the energy-filtered mode in the secondary electrons spectrum range.14 The low-energy cutoff of secondary electrons spectrum (region from 3 to 7 eV, called the emission threshold) is extracted by fitting experimental spectra with a complementary error function model.15 This threshold is directly related to the work function Etr = φsample − φanalyzer, where Etr is the kinetic energy at which the photoemission threshold is observed and φanalyzer is the work function of the
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THEORETICAL SECTION The WF calculations were performed using the QUANTUM ESPRESSO code.19 A local density approximation (LDA) was adopted for the exchange-correlation functional and normconserving pseudopotential for all the chemical elements considered. The kinetic cutoff for the plane-wave basis was 30 Ry. All the atomic positions in the supercell were relaxed until the force acting on each atom was less than 0.003 Ry/au. The Si(100) and Ge(100) systems were modeled by a 37-layers slab where top and bottom layers terminated by 2 H atoms in order to saturate the dangling bonds and to eliminate intragap states. A vacuum layer of more than 20 Å has been used in order to avoid spurious interactions among periodic replicas of the supercell. To simulate the Si1−xGex heterostructure, we modeled an ordered configuration with a variable Ge concentration that preserves the sharp interface between Si and Ge. The slabs thickness of nearly 49−53 Å allows to correctly reproduce the electronic structure of pure bulk Si and Ge (within the LDA accuracy). For all the considered geometries, the in-plane lattice parameter (e.g., in the x and y directions) was that of bulk Si (e.g., aSi Bulk = 5.40 Å) along the xy directions. This means that we have supposed that the Ge layers were grown pseudomorphically on top of pure Si substrates. Clearly the model used is different from the case of a random Si1−xGex film. However, for random alloyed Si1−xGex 26778
DOI: 10.1021/acs.jpcc.5b09278 J. Phys. Chem. C 2015, 119, 26776−26782
Article
The Journal of Physical Chemistry C
Figure 2. XPEEM measurements on the Si1−xGex sample. (a) WF map before sputtering. (b) WF map after Ar+ sputtering. On images b and c, each layer is identified with the same labels than in Figure 1. (c) Averaged 20-line cross sections.
Figure 3. KFM measurements on the Si1−xGex sample. (a) CPD mapping before sputtering, AM mode. (b) CPD mapping after sputtering, AM mode. (c) CPD mapping after sputtering, FM mode. Each layer is identified with respect to the nomenclature defined in Figure 1a. (d) Averaged 20line cross sections are shown.
and averaged 20-line cross sections to obtain WF variations. The maximum scan area allowed by the piezoelectric system is 8 × 8 μm2, which is not enough to map the whole structure. Consequently, two scans (left and right with respect to the glue) are necessary to map the whole region of interest. We notice that, as for XPEEM mapping, WF differences between each layer are visible in KFM images only after sputtering; they were blurred by carbon contamination before. In opposition to spectroscopy, we do not notice a contrast inversion of the images before and after sputtering, but rather a CPD shift to higher values. We assume that carbon contamination is not homogeneous, which explains why images are different before sputtering. Also, we can see that the FM mode provides a better contrast than the AM mode; this is due to the fact that the FM mode signal is proportional to the force gradient and not to the force directly, as in the AM mode.22 Finally, even though WF variations between each SixGe1−x layer are small (
