Defect formation in Si by B10H14 (decaborane) ion implantation has been investigated with photoluminescence measurement. An intense W-line was observed at photon energy of 1.018eV from as-implanted FZ-Si by 30keV B10H14+ implantation. W-line center is considered as an interstitial aggregate and usually observed after ion implantation with subsequent low-temperatureannealing in the case of atomic ion implantation. As W-line is observed from as-implanted Si, the defect formation with B10H14 is expected to be different from that of B+ implantation with the same energy per atom. The energy dependence of W-line intensity is similar to that of diffusivity enhancement after rapid thermal annealing. Molecular dynamics simulation and Rutherford backscattering spectrometry channeling experiment suggest that one B10H14 implantation creates a larger number of dislocated Si atoms than that of B+ implantation with the same energy per atom. This characteristic of B10H14 implantation may cause the different defect reactions in subsequent annealing process.

We have investigated diode leakage in junctions produced by ion-implantation of B with energies of 0.5 - 2 keV and doses of 2 × 1014 — 2 × 1015 cm−2 into n-type wells of ∼1 × 1018 cm−3, after rapid-thermal anneals (RTA) in lamp-based and hot-wall furnaces. Junc- tions are as shallow as 30nm and were directly probed to avoid complications arising from metalization. The leakage current, Ilkg, was found to be independent of the implant dose at any reverse voltage (-1 and -5 V). This implies that the electrically active defects are sufficiently far removed and on the surface-side of the junction. In both systems, a spike anneal (no intentional dwell time at peak-temperature) resulted in higher Ilkg than a soak anneal (dwell time of several seconds at peak-temperature). However, for the same spike annealing recipe, the hot-wall RTA produces tighter distributions than the lamp-based RTA. The width of the distribution is a mea- sure of the temperature uniformity across the wafer. Best leakage currents are of the order 1 × 10−6 A/cm2, in good agreement with device simulations The shallowest junctions exhibit Ilkg ∼5 × 10−4 A/cm2, still well below the specification of even the low power transistor of a 100 nm technology.

Nitrogen diffusion and defect structure were investigated after medium to high dose nitrogen implantation and anneal. 11 keV N2+ was implanted into silicon at doses ranging from 2×1014 to 2×1015 cm−2. The samples were annealed with an RTA system from 750°C to 900°C in a nitrogen atmosphere or at 1000°C in an oxidizing ambient. Nitrogen profiles were obtained by SIMS, and cross-section TEM was done on selected samples. TOF-SIMS was carried out in the oxidized samples. For lower doses, most of the nitrogen diffuses out of silicon into the silicon/oxide interface as expected. For the highest dose, a significant portion of the nitrogen still remains in silicon even after the highest thermal budget. This is attributed to the finite capacity of the silicon/oxide interface to trap nitrogen. When the interface gets saturated by nitrogen atoms, nitrogen in silicon can not escape into the interface. Implant doses above 7×1014 create continuous amorphous layers from the surface. For the 2×1015 case, there is residual amorphous silicon at the surface even after a 750°C 2 min anneal. After the 900°C 2 min anneal, the silicon fully recrystallizes leaving behind stacking faults at the surface and residual end of range damage.

Doping silicon with arsenic by ion implantation above the solid solubility level leads to As clusters and/or precipitates in the form of monoclinic SiAs causing electrical deactivation of the dopant. Information on the local structure around the As atom, and the As concentration depth profiles is important for the implantation and annealing process in order to reduce the precipitated As and maximize the electrically activated As. In this study, we determined the local As structure and the precipitated versus substituted As for As implants in CZ (001) Si wafers, with implant energies between 20 keV and 100 keV, and implant doses ranging from 1 × 1015/cm2 to 1 × 1018/cm2. The samples were subjected to different thermal annealing conditions. We used secondary ion mass spectrometry (SIMS) and UT- MARLOWE simulations to determine the region where the As-concentration is above the solid solubility level. By x-ray absorption fine structure spectroscopy (XAFS), we probed the structure of the local environment around As. XAFS being capable of probing the short-range order in crystalline and amorphous materials provides information on the number, distance and chemical identity of the neighbors of the main absorbing atom. Using Fourier analysis, the coordination numbers (N) and the nearest-neighbor distances (R) to As atoms in the first shell were extracted from the XAFS data. When As precipitates as monoclinic SiAs, the nearest-neighbor distances and coordination numbers are ∼2.37 Å and ∼3, as opposed to ∼2.40 Å and ∼4 when As is substitutional. Based on this information, the critical implant dose where the precipitation/clustering of As starts, and the ratio of the substitutional versus cluster/precipitate form As in the samples were determined.

A model for nucleation and growth of {311} defects is proposed on the basis of thermodynamic and kinetic considerations. Simulated results are discussed and compared to experimental results found in the literature. According to our model it is found that formation energies of self-interstitial clusters depends on the local interstitial supersaturation. Physical parameters extracted from experimental results by inverse modeling are in good agreement with recent values published in the literature.

In this work, the time evolution of B transient enhanced diffusion (TED) suppression due to the incorporation of 0.018% substitutional carbon in silicon was studied. The combination of having low C concentrations, which reduce B TED without completely eliminating it, and having diffused B profiles for several times at a single temperature provides much data upon which various models for the suppression of B TED can be tested. Recent work in the literature has indicated that the suppression of B TED in C-rich Si is caused by non-equilibrium Si point defect concentrations, specifically the undersaturation of Si self-interstitials, that result from the coupled out-diffusion of carbon interstitials via the kick-out and Frank-Turnbull reactions. Attempts to model our data with these two reactions revealed that the time evolved diffusion behavior of B was not accurately simulated and that an additional reaction that further reduces the Si self-inter- stitial concentration was necessary. In this work, we incorporate a carbon interstitial, carbon substitutional (CiCs) pairing mechanism into a comprehensive model that includes the C kick-out reaction, C Frank-Turnbull reaction, {311} defects, and boron interstitial clusters (BICs) and demonstrate that this model successfully simulates C suppression of B TED at 750 °C for anneal times ranging from 10 s to 60 min.

Initial studies (using Scanning Spreading Resistance Microscopy) on the lateral diffusion of B and As have shown an important influence of the thickness of oxy/nitride spacers. The latter phenomenon was tentatively ascribed to stress enhanced diffusion under the spacer region [1]. These studies have been complemented with Scanning Capacitance Microscopy (SCM) measurements, which confirm the SSRM-data. In fact both techniques shows a similar increase in lateral diffusion with increasing spacer thickness (∼ 0.2 nm/nm spacer thickness), whereby no effect is observed on the vertical diffusion. When using spacers with or without TEOS-liner, fairly similar enhancements could be seen. Micro-Raman and CBED stress measurements for these cases do however show a large reduction in stress when a TEOS-liner is used, suggesting that the correlation (at least to the final) stress is not really justified. A possible explanation could however be that the lateral diffusion occurs before the stress relaxation within the thermal treatment. In order to elucidate the diffusion mechanism (initial stress, interstitials, hydrogen incorporation, TED,..) we have expandedthe experimental matrix with a vacancy diffuser such as Sb and simulated the potential H-incorporation duringthe nitride deposition by a hydrogen anneal. Moreover we also have studied the impact of TED by splits with RTP-anneals before the nitride deposition.

Silicon samples were gold-diffused at different temperatures, implanted with He ions at 1.6 MeVand then annealed at 1050°C for 2 hours. The implantation induced-defect structure and their distributionin the depth of the sample, studied by conventional and high resolution cross section electron microscopy (HRXTEM) depend on the gold level introduced in the wafer prior to the gettering process. A high concentration of gold in silicon seems to influence the defect configuration in the cavity zone. Indeed, gold chemisorbed atcavities can homogenize the surface energy of their planes in different orientations, and can increase the cavity critical diameter beyond they become facetted. Secondary ion mass spectroscopy (SIMS) profiles exhibit ashouldered shape and a width closely related to the presence of the defects (observed by XTEM) which are veryefficient sinks both for gold and copper atoms. Unfortunately, the electrical improvement of the material (checked by minority carriers diffusion length measurements MCDL) is not achieved by this gettering process, probably due to the high metal impurity concentrations remaining out of the gettering zone, to the presence of AuCu complexes and η-Cu3Si precipitates identified by deep level transient spectroscopy (DLTS)measurements and HRXTEM observations respectively.