Roles of Impurities and Implantation Depth on He + - Cavity Shape in Silicon

Transcription

1 MRS Proceedings; volume 864 / 2005 Roles of Impurities and Implantation Depth on He + - Cavity Shape in Silicon Gabrielle Regula, Rachid El Bouayadi, Maryse Lancin, Esidor Ntsoenzok 1, Bernard Pichaud, Marie-Odile Ruault 2 Laboratoire TECSEN CNRS UMR-6122, Aix-Marseille III, Service 151, Marseille, F CERI-CNRS, 3A, rue de la Férollerie, Orléans cedex, F CSNSM, CNRS-IN2P3, bâtiment 108, Orsay, F ABSTRACT Silicon samples were implanted with He + ions at energies varying from 10keV to 1.55MeV using doses ranging from cm -2 to cm -2 to obtain similar He concentration at each projection range (R p ). In few samples, gold, platinum, nickel or silver was introduced prior to He + implantation by diffusion at temperatures ranging from 870 C to 1050 C. All samples were annealed in the 400 C 1050 C temperature range to determine the equilibrium stage of the growth of the cavity. The cavity characteristics (distribution, shape and size) were studied by cross section transmission electron microscopy (XTEM). Their morphology demonstrates the validity of the chemisorption hypothesis when they grow in silicon intentionally contaminated by metal. A consequence of the surface proximity on the cavity characteristics was verified and allows stepping forward two regimes of cavity growth: one, very fast, taking place in a He-free environment and another one, slower, occurring in a He-rich atmosphere. INTRODUCTION Cavities can be introduced in silicon by light gas ion (like He + ) implantation if a minimum concentration of at/cm 3 is reached [1]. They are being studied for more than two decades since they are very interesting from both fundamental and applied-research. As for the first point of view, a complete surface energy plot of silicon can be obtained by studying the equilibrium shape of the cavity reached after appropriate thermal budget, using both Wulff constructions and high resolution microscopic observations [2]. The binding energy of different impurities [3] as well as data characteristic of gas out-diffusion from an ideal surface, [4,5] can also be derived from secondary ion mass spectrometry (SIMS) or Rutherford backscattering spectrometry (RBS) of cavities. The nucleation of the cavities was studied theoretically. The simulations showed [6] that He atoms settle preferentially in an interstitial tetrahedral site (T s ) and are not stable in a single vacancy (V). More likely, He atoms agglomerate in adjacent T s inducing a local compressive strain that can be compensated by the formation of He m -V n complexes [7]. He-V 2 complexes are assumed to be the earliest precursors of the cavity formation [8,9]. The V surrounding of He atoms induces a tensile strain which acts as a driving force to drag He, building an onion-like structure. The He m -V n complexes form in the implanted-zone rich of both He and V [10] and their dissociation occurs at about 250 C [11]. Concerning the cavity growth, two mechanisms were proposed. Some authors ascribe the growth to the migration of small cavities to bigger ones until they coalesce [12,13], though other explain it by exchanges of entities (V and He) that is to say by an Oswald ripening [13,14]. The latter

2 mechanism must be applied to a conservative system and thus should be greatly limited by the loss of V at the sample surface. According to nuclear reaction analysis (NRA), the desorption rate decreases significantly in the early stage of annealing [15,16]. These results suggest that He atoms remain in the sample even after high temperature annealing. The parameters which can control the cavity growth are numerous: the local concentrations of V and self-interstitial at Rp, the distance between the cavities and the surface and the impurity concentration in the Si matrix [15,7-20]. So far, the nucleation and growth mechanisms of cavities are still unclear and under debate. In addition to the fundamental knowledge, the cavities are very attractive for microelectronic applications. They can be introduced in the devices to control the charge carrier lifetime, to exfoliate thin silicon films, or to create in the devices local gettering sites which allow a reduction of the thermal budget and the trapping of impurities in concentration far lower than their solubility limit. Indeed, the gettering ability of cavities has been demonstrated for light elements [19] as well as for transition metals both at He + implantation energies in the kev [21] and MeV [22] ranges. Depending on both the experimental conditions and the nature of the impurity (interstitial, hybrid, forming silicides or not), the atoms are supposed to be trapped at cavity either as a fraction of a chemisorbed-layer [23], up to a monolayer (which is too thin to be detected by microscopic studies) or as a three dimensional structure [24]. Moreover, the gettering efficiency of the cavities was found to be dependent of the cavity radius [25]. However, the chemisorption hypothesis [23] was never confirmed by direct microscopic observation. The first part of this paper deals with the determination of the equilibrium shape of the cavities induced by an implantation of He + cm -2, at 1.55 MeV, for different annealing temperatures up to 1050 C. In the second and third part, the determination of the cavity shape is used as a tool to check the chemisorption mechanism and the possible role of He in the cavity growth. EXPERIMENTAL DETAILS Three kinds of (111) silicon, 500 µm thick, were studied: they were grown either by epitaxy (EG), or by float zone (FZ) or by Czochralski (Cz) methods. EG samples consist of a n-type (30 cm) layer, 75 µm thick, on a n-type Czochralski substrate (5m.cm). FZ and Cz samples are p type, 6-10.cm and 7-8.cm respectively. He + implantations were performed in the kev energy range with an implanter equipped with an Abernas-Niels source-type and at 1.55 MeV via a van de Graff accelerator. All implantations were carried out using different He + doses (see table I) to get similar maximum He + concentration at the R p predicted by transport range of ions in matter (TRIM). All thermal treatments were performed under N 2 or Ar gas flow in a quartz tube in a conventional furnace, and then cooled to room temperature with a rate of about -5 C.s -1 for the first 500 C loss. The induced-cavity growth was studied after one hour annealing at temperatures from 400 C to 1050 C. The foils prepared for XTEM studies were either prepared by conventional mechanical thinning and ion milling or cut via a focused-gallium ion beam. Energy (kev) Table I: Implantation parameters He + dose (cm -2 ) used to compare the growth when Rp (nm) the cavities are created either close max ([He]) (cm -3 ) or far from the sample surface.

3 For chemisorption studies, some samples were voluntary and uniformly contaminated by high level of metal (table II) prior to the MeV implantation. Thus, the cavity growth was performed at 1050 C for two hours in an impurity rich environment. To determine the global cavity coverage [18] SIMS depth profiles were carried out with a Cs + or O + primary source (CAMECA IMS4F) and XTEM images were performed (Jeol 2010F and Philips CM12). When precipitates were observed, their composition was verified by energy dispersive X-ray spectrometry (EDS). Table II: Description of the processes developed to verify the chemisorption hypothesis by XTEM observations of the equilibrium shape modification of the cavity. The specific surface energies of metals are also given. Metal/ diffusion silicide Metal T diffusion He + = cm -2 T growth sample mechanism formation source ( C) Energy (MeV) ( C) (Jm -2 ) Au/ FZ hybrid no Au-Si eutectic Ni/Cz interstitial yes Film evaporation Ag/Cz interstitial no conventional conducting paint Pt/EG hybrid yes PtSi film RESULTS AND DISCUSSION Bulk cavity growth Figure 1 shows XTEM micrographs of six different samples implanted at 1.55 MeV and annealed at temperatures from 400 C to 1050 C. There is an expected-evolution of the spatial distribution size and morphology of the cavities since the studied-system is assumed to be conservative. Figure 1. Bright field XTEM observations along a <110> of EG and FZ samples implanted with He + at 1.55 MeV/ cm -2 and annealed one hour at a) 400 C; b) 700 C; c) 800 C; d) 850 C; e) 900 C: note that the biggest cavities are hardly surrounded by smaller ones; f) 2 hours at 1050 C. The surface location corresponds to the upper part of each picture.

4 As the temperature increases, the width of the cavity band shrinks from 400nm till forming a 300nm-wide chain in planes parallel to the sample surface. The mean cavity radius increases from 4nm up to 100nm while the cavity density drops from µm -2 to µm -2. The cavity radius distribution is homogeneous after a low temperature annealing. For higher temperature up to 900 C, the scattering of the cavity radius increases, the biggest ones corresponding to cavities located nearby the middle of the cavity band thickness. It is noteworthy that the larger the cavities, the lower the density of small cavities in their vicinity. At the highest annealing temperature (1050 C, 2 hours), the scattering of the cavity radius decreases strongly, and all cavities are found to be facetted. At lower annealing temperature it is very difficult to know what is up for faceting of the smallest cavity. A further 2 hours annealing at 1050 C gives the same result as displayed in figure 1.f. Thus, we assume that the cavities have reached their equilibrium distribution after a single thermal treatment of 2 hours at 1050 C and form a quasi monolayer of facetted-cavity. The corresponding cavity morphology, close to the one described by Eaglesham et al. [2] will be taken as reference in the next part. Chemisorption validation SIMS measurements (not shown) demonstrate that all types of metallic impurities (Au, Pt, Ni, Ag) can be trapped by the cavities even in the presence of residual He pressure (about 20% of the initial He + dose [15]). The metal coverage can be estimated by coupling SIMS analyses with XTEM observations [18]. Figure 2 shows the striking morphology evolution of big cavities grown in different metal-rich atmospheres during an annealing for 2 hours at 1050 C, provided is higher than 1%. It was demonstrated that the residual He pressure does not influence the shape of cavities the diameter of which is larger than 10 nm [18]. Therefore, the cavity morphology can be explained only in terms of specific surface energy, the elastic energy being neglected. In the range of 1%-10%, values associated to clean Si vicinal surfaces, which are assumed to contain the preferential trapping sites, is decreased down to the Si lowest value {111} =1.23 Jm -2 [2] whatever the metal trapped [26]. The consequence is the disappearance of facets even for the biggest cavities. Figure 2. Bright field XTEM observations along <110> of EG and FZ samples metal-diffused at 1050 C for 2 hours, implanted with He + at 1.55MeV/ cm -2 and annealed a second time at 1050 C for 2 hours: evolution of cavity morphology with the metal coverage. Close to a chemisorbed-monolayer, of the cavity vicinal surfaces become equal to that of the adsorbed-metal. For instance, concerning Pt the specific surface energy of which is much higher

5 than those of vicinal facets of Si, the {111} faceting is strongly favored when the coverage is close to a monolayer. The cavity shape must depend both on the three dimensional phase grown in the cavity and the cooling conditions when value exceeds 100%. As for Ag, it precipitates inside cavities with an equilibrium shape very similar to the one of clean cavity. A question remains though: since He does not step in metal trapping or cavity morphology, does it play any role in cavity formation? To answer this question, the next section reports studies of cavity growth kinetics after implantation in the kev energy range. Near-surface cavity growth: role of He The quasi monolayer of facetted-cavity is obtained after a 900 C annealing as shown in figure 3a. Thus, a lower thermal budget for kev implantations than for MeV ones (1050 C) is needed to reach what is assumed to be the last stage of cavity growth. Note that with the same thermal budget (900 C), both energy range implantations do not result in the same spatial distribution of cavities. Indeed, small cavities cannot be observed in figure 3a whereas they are still numerous in figure 1e. Such difference strongly suggests that the cavity growth is slowed down by the presence of He deep in the bulk, in agreement with Grisolia et al [27]. Indeed, they reported that He diffusion was the slowest process (with an activation energy of 1.70 ev). In the frame of an Oswald ripening mechanism, it is assumed that in rich-he environment, cavities grow exchanging both He and V. In He-free silicon (which is the case for low implantation energy), cavities grow faster since they exchange only V. Two regimes of cavity growth can then be defined in terms of kinetic process depending on the He level. Figure 3. Bright field XTEM observations along <110> of EG samples annealed at 900 C after implantation step at a) 50 kev, He + cm -2 and b) 10 kev, He + cm -2. The surface location corresponds to the upper part of each picture. Nevertheless, some points concerning the growth mechanism of cavities remain unclear and the He regimes has to be precised. Moreover, how calling the mechanism if cavities do not exchange point defects but small He m V n complexes (likely HeV 2 ) during the thermal treatment. The cavity growth process would be a hybrid of the Oswald ripening and of the coalescence mechanism. To emphasize the limits of the model, let us focus on the cavity shape when the cavity layer is very close to the surface. Comparison of figure 3a and 3b shows that the closer the cavity, the less facetted. This result is either the consequence of chemisorption due to a slight surface contamination or more likely, it means that the surface itself plays a role in the cavity growth process. CONCLUSIONS The chemisorption is proofed both chemically by SIMS and physically by XTEM pointing out a striking change in the equilibrium shape of cavities with the nature and amount of impurity present in the wafer during the cavity growth. He + was not found to prevent chemisorption but to slacken cavity growth (comparing kev and MeV implantations). In He-rich environment (in bulk

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