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Impact of alloying elements (Co, Pt) on nickel stanogermanide formation
Materials Science in Semiconductor Processing 108 (2020) 104890
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Impact of alloying elements (Co, Pt) on nickel stanogermanide formation Andrea Quintero a,b , Patrice Gergaud a , Jean-Michel Hartmann a , Vincent Reboud a , Eric Cassan b , Philippe Rodriguez a ,∗ a b
Univ. Grenoble Alpes, CEA, LETI, F-38000 Grenoble, France Centre de Nanosciences et de Nanotechnologies, CNRS, Univ. Paris-Sud, Université Paris-Saclay, C2N - Orsay, 91405 Cedex, France
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Keywords: Advanced metallization Ni NiCo NiPt Solid-state reaction Stanogermanide GeSn
ABSTRACT The impact of Pt or Co as alloying elements for Ni-based metallization of GeSn layers has been investigated. As far as the solid-state reaction is concerned, the overall phase sequence is the same for all metallizations: at low temperature, a Ni-rich phase is obtained; it is then consumed to form the low resistivity mono-stanogermanide phase. Nevertheless, the addition of an alloying element has an impact on Ni consumption, Ni-rich and monostanogermanide phases’ formation temperatures. Moreover, the addition of Co or Pt positively impacts Sn segregation by delaying this phenomenon. Co has a weak influence on morphological and electrical properties. On the other hand, Pt improves the surface morphology by delaying the Ni(GeSn) phase agglomeration and enhancing the process window in which the sheet resistance remains low.
1. Introduction Ge1-x Snx alloys have seen their interest growing in the past few years. These materials are envisioned for electronic applications such as source and drain stressors in Ge MOSFETs [1] and high mobility channels in pMOSFETs [2] or pTFETs [3]. Since the addition of enough Sn in the Ge lattice structure (about 10 at.%) results in a direct band-gap semiconductor, Ge1-x Snx layers have also been considered for optoelectronic applications such as p-i-n photodetectors [4,5], heterojunction LEDs [6,7] and, more recently, optically pumped lasers [8,9]. Whatever the targeted application, an efficient electrical ohmic contact is needed to obtain high performance devices. Ni-based contacts have been proposed and widely studied as they enable the formation of Ni(GeSn) intermetallics at relative low temperature with low values of contact (Rc ) and sheet (Rsh ) resistances [10, 11]. Nevertheless, Ni-based contacts exhibit a poor thermal stability linked to the precipitation of Sn atoms and the agglomeration of the Ni(GeSn) layer. These phenomena are enhanced by the increase of Sn content in the Ge1-x Snx layers [12,13]. To extend the thermal stability of Ni-based contacts, solutions have been proposed such as the cosputtering of Ni and Pt [14,15] or the addition of a Pt interlayer [16, 17]. As an alternative to Ni-based contact and in order to extend the thermal stability of contacts, Ti-based contacts have also been evaluated [18–20]. In this study, the impact of alloying elements, namely Co and Pt, on the Ni/GeSn system was investigated. The addition of Pt or Co should enhance the thermal stability of stanogermanides, given what is known
for the Ni/Si(Ge) systems [21–26]. Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 metallizations on Ge0.9 Sn0.1 layers were thus investigated. In particular, the impact of alloying elements on the solid-state reaction and the contact properties – in terms of surface morphology and electrical properties – was studied. The aim of this study was to highlight the opportunities of using alloying elements and CMOS-compatible metallization schemes to enhance the thermal stability of Ni-GeSn contacts. 2. Experimental GeSn layers with 10 at.% of Sn were epitaxially grown at 325 ◦ C, 100 Torr with Ge2 H6 and SnCl4 precursors on Ge-buffered Si (1 0 0) substrates in reduced pressure chemical vapor deposition (RPCVD) chambers [27]. The Sn composition was compatible with the optoelectronics applications described in the introduction. The thickness of the GeSn layers (60 nm) was low enough for them to be fully compressively strained on the Ge strain relaxed buffers (SRB) underneath. The layer thickness, 60 nm, and the Sn content, 10 at.%, were extracted from 𝜔–2𝜃 scans around the (0 0 4) XRD order. Surface preparation before metal deposition consisted in dips for 1 min in hydrofluoric acid (HF) diluted at 1% in water followed by rinses in deionized water and drying with a N2 gun – in order to remove particles and native oxides from the surface. Then, the deposition, at room temperature, of 10 nm of Ni, Ni0.9 Co0.1 or Ni0.9 Pt0.1 capped with 7 nm of TiN was performed by magnetron sputtering in a 300 mm cluster tool. To limit the regrowth of native oxides, the time interval between the surface preparation and
∗ Corresponding author. E-mail address: [email protected] (P. Rodriguez).
https://doi.org/10.1016/j.mssp.2019.104890 Received 11 September 2019; Received in revised form 18 November 2019; Accepted 16 December 2019 1369-8001/© 2019 Elsevier Ltd. All rights reserved.
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Fig. 2. In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni0.9 Co0.1 and Ge0.9 Sn0.1 .
Fig. 1. In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni and Ge0.9 Sn0.1 .
fluctuation of the hexagonal metastable 𝜖-Ni5 (GeSn)3 phase when the Ge(Sn) composition varied [28]. At 120 ◦ C, Ni was totally consumed. At 210 ◦ C, the diffraction peaks linked to the Ni-rich phase dissipate and new diffraction peaks appear at 42.4◦ , 45.3◦ and 53.4◦ . They correspond to the formation of the Ni mono-stanogermanide phase Ni(GeSn) with the following orientations: (2 1 0), (2 1 1) and (0 1 3). This phase remains stable until 675 ◦ C, a temperature at which the diffraction peaks vanish. This is due to a phase destabilization when reaching the NiGeSn–GeSn eutectic point. Even if the eutectic point between NiGe and Ge was reported to be at around 760 ◦ C [29], the addition of Sn in the Ge lattice might modify this temperature. In addition, as for all annealing equipment, the temperature indicated by the furnace and the absolute temperature might differ by a few tens of degrees. In order to study a potential impact of the addition of 10 at.% of Co on Ni-based metallization, the Ni0.9 Co0.1 /Ge0.9 Sn0.1 solid-state reaction was monitored by in-situ XRD. Patterns are shown in Fig. 2. The phase sequence is the same than for the Ni/GeSn system: the Ni0.9 Co0.1 (diffraction peak at 44.4◦ ) is consumed to the benefit of a NiCo-rich phase (NiCo)5 (GeSn)3 (diffraction peak at 46.5◦ ) which is in turn consumed to form the NiCo mono-stanogermanide phase (NiCo)(GeSn) (diffraction peak at 2𝜃 = 53.9◦ ). Nevertheless, a few discrepancies are seen. First, the NiCo-rich phase is present in films right after the deposition state. At high temperature, after the disappearance of the (NiCo)(GeSn) phase when reaching the NiGeSnGeSn eutectic point, another diffraction peak at 34.7◦ is observed. This diffraction peak was not detected during comparable in-situ XRD measurements on Ni/GeSn and NiCo/Ge systems. This peak is thus due to the concomitant presence of Co and Sn and the formation of CoSnx compounds. The impact of the addition of 10 at.% of Pt on the Ni-based solidstate reaction was also studied by in-situ XRD. Fig. 3 shows the contour map obtained for the Ni0.9 Pt0.1 /Ge0.9 Sn0.1 system. This system is by far the most complicated one among the three systems investigated in this study. A complete analysis of the Ni0.9 Pt0.1 / Ge0.9 Sn0.1 system is reported elsewhere [30]. At relative low temperatures, the phase sequence is similar to the one obtained with Ni or Ni0.9 Co0.1 : after Ni0.9 Pt0.1 consumption, the (NiPt)5 (GeSn)3 rich phase is formed. The later is consumed to form the NiPt monostanogermanide phase (NiPt)(GeSn). At around 350 ◦ C, a shift towards higher 2𝜃 angles is observed. This shift appears concomitantly with the appearance of new diffraction peaks at 39.2◦ and 42.0◦ . This is due to a rejection of Pt from the (NiPt)(GeSn) lattice and a transition towards Ni(GeSn) with the formation of PtSnx compounds with various stoichiometries and thermal stabilities [30].
the metallization of the samples was reduced to the minimum (less than 10 min). In addition, there was no air break between Ni(Co,Pt) and TiN capping layers’ deposition: samples were indeed kept at pressures less than 10−7 Torr during transfer between one chamber and another. The TiN layers preserved the Ni(Co,Pt) layers from atmospheric contamination. This way, solid-state reaction between Ni(Co,Pt) and Ge0.9 Sn0.1 was not impacted by oxygen or carbon contamination. The Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 phase formation sequence and crystalline evolution were followed by in-situ X-ray diffraction (XRD) using an Empyrean PANalytical X-ray diffractometer equipped with a copper (Cu K𝛼) source, a linear detector (PIXcel1D ) and a HTK 1200 Anton Paar furnace under secondary vacuum (pressure was maintained at about 10−6 bar to avoid atmospheric contamination). 𝜃–2𝜃 scans over a large range of diffraction angles (20-90◦ ) were performed with a 2◦ offset in 𝜔 in order to attenuate the symmetric (4 0 0) substrate reflection. Profiles were acquired from 50 up to 600 or 700 ◦ C, with 5 ◦ C steps and a scanning time of 7 min per temperature step. Phase analyses were performed with the HighScore Plus software from PANalytical, using the International Centre for Diffraction Data (ICDD) PDF-4 database. Furthermore, samples metallized with Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 using the same procedure as described above were ex-situ annealed for 30 s at temperatures ranging from 150 up to 550 ◦ C in a rapid thermal annealing (RTA) furnace with a N2 environment. Then, 𝜃–2𝜃 scans over a large angular range (20–90◦ ) using a Cu source were performed in a X’Pert Pro PANalytical X-ray diffractometer. The surface morphology and the electrical characteristics were assessed by tapping mode atomic force microscopy (AFM) with a Bruker FastScan equipment and sheet resistance (Rsh ) with a four-point probe meter. 3. Results & discussion 3.1. Study of the solid-state reaction The Ni/Ge0.9 Sn0.1 solid-state reaction followed by in-situ XRD is shown in Fig. 1 as a contour map. The 𝑦-axis shows the temperature and the 𝑥-axis, the diffraction angles. Peaks of TiN (1 1 1) at 36.4◦ , Ni (1 1 1) at 44.9◦ , Si (4 0 0) and Ge/GeSn (4 0 0) around 68◦ are observed. At low temperature (above 85 ◦ C), new diffraction peaks appear at 2𝜃 around 46.6◦ and 85.6◦ . These peaks are due to the formation of the 𝜖-Ni5 (GeSn)3 rich phase [28] with the following orientations: (1 1 0) and (3 0 0). The shift towards higher 2𝜃 angles observed at 130–140 ◦ C was linked to stoichiometry variations, i.e. lattice constant 2
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Fig. 5. Phase sequences of the reaction of 30 nm-thick Ni, Co or Pt films on Ge (0 0 1). Source: Adapted from [32].
Fig. 3. In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni0.9 Pt0.1 and Ge0.9 Sn0.1 .
soluble in Ni5 Ge3 or NiGe should decrease the nucleation barrier [31]. The phase sequences of Ni, Co and Pt germanides are depicted in Fig. 5. The NiGe phase can coexist with (i) CoGe, Co5 Ge7 and CoGe2 phases and (ii) Pt2 Ge, Pt3 Ge2 , PtGe, Pt2 Ge3 and PtGe2 phases [32]. NiGe being orthorhombic, all the orthorhombic phases might be soluble in NiGe. For the Pt/Ge system, all the intermetallic phases, except Pt2 Ge which is hexagonal, are orthorhombic. For the Co/Ge system, CoGe2 only is orthorhombic. Indeed, CoGe is monoclinic and Co5 Ge7 tetragonal. Thus, based only on solubility point of view, the fact that Co decreases the formation temperature of NiGe whereas Pt shifts it up towards higher temperatures cannot readily be explained. The Ni-rich phase, Ni5 Ge3 , coexists only with transition metals. The fact that the Co has an hexagonal structure, like Ni5 Ge3 while Pt has a cubic structure might explain why the addition of Co decreases the energy threshold for the Ni-rich phase. For the mono-germanide phase, if the impact of the alloying element on the nucleation barrier is not obvious, one should consider surface/interface effects. The Co–Ge, Ni–Ge and Pt–Ge bond strengths are equal to 230, 290 and 400–500 kJ mol−1 , respectively [33–35]. If the bonding between atoms is weaker (which is the case for Co), surface, interface and grain boundary diffusions are easier. As a consequence, atomic diffusion will be faster, resulting in a shift of the formation of intermetallics towards lower temperature with Co. The same reasoning would explain why the addition of Pt shifts the formation of intermetallics towards higher temperatures. In addition, it can be seen in Fig. 3 that, at around 120 ◦ C, the Ni0.9 Pt0.1 alloy peak shifts towards lower 2𝜃 angles, suggesting an increase of the alloy lattice parameter. The phenomenon is known as Pt snowplow and has already been discussed elsewhere [30]. This phenomenon might explain the delay noticed for the appearance of the Ni-rich phase with the addition of Pt.
Fig. 4. Impact of Co or Pt addition on Ni consumption, Ni-rich and monostanogermanide phases’ formation temperatures.
The study of Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 solid-state reactions has highlighted that, if the overall phase sequence is the same, discrepancies do exist. Fig. 4 summarizes the impact of Co or Pt addition on Ni consumption, Ni-rich and mono-stanogermanide phases’ formation temperatures. As far as Ni consumption is concerned, the addition of 10 at.% of Co does not impact the temperature at which the Ni is totally consumed. Indeed, Ni full consumption is observed at 120 ◦ C for both metallizations. On the other hand, the addition of 10 at.% of Pt pushes the Ni total consumption up to 140 ◦ C. Pt redistribution during thermal reaction, in particular the Pt snowplow effect, might explain this delay in total Ni consumption [30]. The formation of the Ni-rich phase is strongly impacted by the nature of the alloying element. For Ni metallization, the Ni5 (GeSn)3 phase appears at 85 ◦ C and is totally consumed at 210 ◦ C. When Co is added, the (NiCo)5 (GeSn)3 phase is observed since the as-deposited state. This phase is totally consumed at 210 ◦ C. On the other hand, the addition of Pt postpones the formation of the Ni-rich phase to 140 ◦ C and its total consumption to 240 ◦ C. Similar trends are observed for the temperature at which the mono-stanogermanide phase appears. Co addition tends to lower this temperature (200 ◦ C vs. 210 ◦ C for Ni) whereas Pt addition pushes this temperature up to 240 ◦ C. To sum up, the addition of Co lowers the formation temperatures of the Ni-rich and mono-stanogermanide phases whereas Pt addition delays the formation of these intermetallic phases. The impact of the alloying element’s nature on the formation temperature of intermetallics can be discussed according to various lines. From a solubility point of view, elements forming germanides that are
3.2. Impact on the Sn segregation There was, upon Co or Pt addition, a formation of CoSnx and PtSnx compounds at high temperatures (see the solid-state reactions described in Section 3.1). The formation of these kinds of compounds suggests the presence of Sn in the metallic state. However, no diffraction peak was linked to the Sn phase in Figs. 1 to 3. Indeed, beyond 230 ◦ C, tin is in a melted state and no diffraction peak can be detected with in-situ XRD analyses. Thus, XRD analyses were systematically performed on samples annealed by RTA to highlight the impact of the alloying element on the Sn segregation threshold. The Sn segregation threshold temperatures for each system, Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 metallizations on Ge0.9 Sn0.1 , are plotted in Fig. 6. As far as the Ni/Ge0.9 Sn0.1 system is concerned, no diffraction peak linked to the 𝛽-Sn phase was discernible at 300 ◦ C. Diffraction peaks associated with the 𝛽-Sn phase were however observed since 350 ◦ C. Thus, the Sn segregation threshold for this system is ranging between 300 and 350 ◦ C (hatched area in Fig. 6). In the same way, the 3
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Whatever the system studied, at low temperature (i.e. for T ≤ 300 ◦ C), samples exhibit cross hatch patterns along the ⟨1 1 0⟩ crystallographic direction. This pattern is due to the propagation of the 60◦ threading arms of misfit dislocations on (1 1 1) planes during Ge buffer and subsequent GeSn epitaxial growth [27,42]. This means that, in the as-deposited state, Ni-based and TiN layers are extremely smooth and their deposition does not modify the original surface morphology. This morphology is retained up to 300 ◦ C, meaning that the formation of the Ni-rich phase does not modify the original morphology. At 350 ◦ C, the surface morphology for Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems drastically changes, with a strong increase of the surface roughness. This change has been linked to the formation of the monostanogermanide phase and its subsequent agglomeration [30]. At higher temperatures, the roughness gradually increases due to the amplification of both phenomena: the mono-stanogermanide phase agglomeration and Sn segregation. The surface morphology evolution is very similar for Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems. Thus, the addition of Co as an alloying element does not really impact the surface morphology. The analysis of the Ni0.9 Pt0.1 /Ge0.9 Sn0.1 system highlights substantial differences in terms of surface morphologies. At 350 and 400 ◦ C, even if the morphology is changing with the appearance of cracks at the sub-micron scale on the surface, the cross-hatch is still visible. These cracks are linked to grain boundary grooving phenomenon at the early stage of agglomeration. The cross-hatch is no longer observable at 450 ◦ C and above, and the surface roughens. Fig. 7 clearly shows that the addition of Pt has a positive impact on surface morphology, by delaying the agglomeration of the Ni(GeSn) layer. An in-depth discussion can be found in a previous paper [30]. Briefly, the Pt–Ge bond energy being higher than the Ni–Ge one, Pt atoms will reduce the breakdown rates of chemical bonds which will reduce grooving and could explain the morphological stability improvement at high temperatures. Pt addition might also suppress or limit Ge out diffusion, which is beneficial for the morphological stability [35]. It has also been demonstrated that Pt addition reduces the grain size of the monostanogermanide phase [30], reducing agglomeration phenomena [43, 44].
Fig. 6. Sn segregation threshold extracted from XRD analyses of samples annealed by RTA.
Sn segregation threshold for the Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems are ranging between 350 and 400 ◦ C and 400 and 450 ◦ C, respectively. Thus, the addition of Co or Pt increases the Sn segregation temperature. The GeSn layers being grown by RP-CVD under kinetic conditions far from equilibrium, they are metastable. Indeed, GeSn is known for its poor thermal stability. Sn has been reported to segregate at the surface when the annealing temperature increases [36,37]. Sn segregation is linked to the formation of Sn clusters/precipitates in the GeSn layers and the subsequent migration of atoms towards the surface [36–38]. It can occur more easily near the surface because the surface free energy is higher than the bulk free energy [36]. Even if reducing the density of point defects limits the initial formation of Sn clusters/precipitates, Sn segregation is mainly governed by the migration/diffusion of atoms. Thus, the addition of alloying elements might impact (i) the diffusion of Sn atoms and (ii) the surface/interface energies involved in Sn segregation mechanisms. As reported in Section 3.1, the addition of Co or Pt leads to the formation of CoSnx and PtSnx compounds. On the other hand, the formation of NiSnx compounds was not observed. After the formation of Sn clusters/precipitates in GeSn layers, Sn atoms diffuse towards the surface. This diffusion can occur within the grains or, much more likely, at the grain boundaries of Ni(GeSn). In the presence of alloying elements, such as Co or Pt, the Sn diffusion might be slowed down by the formation of CoSnx and PtSnx compounds, resulting in a delayed Sn surface segregation. The second hypothesis concerns the modification of the surface/ interface energies by the alloying element. It is well known that the addition of alloying elements can impact the texture of silicides and germanides [39]. As seen in Figs. 1 to 3, the number of out-of-plane orientations for the mono-stanogermanide phase decreases with the addition of an alloying element: from 3 with pure Ni to only one orientation with the addition of Co and Pt. The formation of highly textured films has already been reported for Pt addition in Ni-based silicides [21,40]. They are expected to have a significantly lower interfacial energy compared to polycrystalline films. Thus, the formation of more textured mono-stanogermanide films with the addition of an alloying element can result in lower the surface/interface energies and have a positive impact on Sn segregation. Finally, Sn segregation has also been linked to strain at the interface [41]. The covalent radius of Ni (124 pm) being lower than the ones of Co (126 pm) and Pt (136 pm), the addition of these elements might modify the local strain and positively impacts the Sn segregation threshold. As seen in Fig. 6, the higher the element’s radius is, the higher the Sn segregation threshold is.
3.4. Impact on the electrical properties Fig. 8 shows the evolution of the sheet resistance after various RTA from 200 up to 550 ◦ C for the Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems. Whatever the system, from the as-deposited state and up to 300 ◦ C, sheet resistances (Rsh ) remain constant at around 28–35 Ω/sq. This temperature range corresponds to the Ni-rich phase domain. At 350 ◦ C, minimum values are reached for all systems: 23.1, 16.3 and 10.4 Ω/sq for the Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems, respectively. Differences are significant. The mono-stanogermanide layer obtained at 350 ◦ C with the addition of Pt as an alloying element is less resistive than the one obtained with Co and even less resistive than the one without any alloying element (pure Ni). These differences might be due to thickness variations, modifications of the intrinsic properties of the intermetallic layers by the alloying element or to the layer’s morphology itself. For Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems, a strong increase of the Rsh values is observed above 350 ◦ C. This behavior is directly linked to the agglomeration of the NiGeSn layer and to Sn segregation. Thus, as observed for the morphology, the addition of Co does not significantly impact the electrical properties. The decrease of the Rsh values at high temperature observed for the Ni0.9 Co0.1 /Ge0.9 Sn0.1 system might be linked to the formation of CoSnx compounds described in Section 3.1. On the other hand, the addition of Pt as an alloying element enhances the process window in which Rsh values remain low. Indeed, a plateau is observed between 350 and 400 ◦ C and the increase is moderate for 450 ◦ C. Beyond 450 ◦ C, the sheet resistance drastically increases, reaching values similar to that of the Ni/Ge0.9 Sn0.1 system.
3.3. Impact on the surface morphology Fig. 7 shows AFM images for Ni (top row), Ni0.9 Co0.1 (middle row) and Ni0.9 Pt0.1 (bottom row)/Ge0.9 Sn0.1 samples annealed at various temperatures. 4
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Fig. 7. Surface AFM images (5 μm x 5 μm scan sizes) of samples annealed at various temperatures. Ni (top row), Ni0.9 Co0.1 (middle row) and Ni0.9 Pt0.1 (bottom row)/Ge0.9 Sn0.1 systems. Scan directions along ⟨1 0 0⟩.
and the sheet resistance evolution as a function of annealing temperature are quite similar for Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems. On the other hand, Pt positively impacts the surface morphology by delaying the Ni(GeSn) phase agglomeration and enhancing the process window in which the Rsh values remain low. Thus, even if both elements have a positive impact on Sn segregation, the impact on morphology, i.e. the ability to delay the agglomeration of the Ni(GeSn) phase, prevails in order to optimize the electrical properties. The addition of Co as an alloying element has a positive impact on Sn segregation, it also decreases the formation temperature of the Ni(GeSn) phase which might be valuable for some applications. Nevertheless, Pt addition, by enhancing the morphological and the electrical properties – mainly by delaying the Ni(GeSn) agglomeration and Sn segregation – is the most interesting option in terms of integration issues. It opens up the way towards a CMOS-compatible contact technology on GeSn. It was used on photonics components such as midinfrared photodiodes and LEDs. First results showed a strong reduction of the contact resistance compared to classical Ti/Al contacts and real improvements in terms of optoelectronic devices performances [45,46].
Fig. 8. Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for the Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems.
Declaration of competing interest 4. Conclusions The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
In this work, the impact of alloying elements, namely Co and Pt, on Ni stanogermanide formation and properties was investigated. As far as the solid-state reaction is concerned, the overall phase sequence is the same: at low temperature, a Ni-rich phase is obtained, the latter being consumed to form the mono-stanogermanide phase. Nevertheless, the addition of an alloying element has an impact on Ni consumption, Ni-rich and mono-stanogermanide phases’ formation temperatures. Based only on solubility point of view, the fact that Co decreases the temperature of formation of NiGe whereas Pt shifts it up towards higher temperatures cannot be readily explained. However, the impact of alloying element on atomic diffusion might affect surface/interface processes. The addition of an alloying element has also a positive impact on the Sn segregation threshold. Indeed, the addition of Co or Pt delays the Sn segregation temperature. The formation of CoSnx or PtSnx might slow Sn diffusion towards the surface and its subsequent segregation. Co and Pt addition might also lower the surface/interface energies, with a positive impact on Sn segregation. If both elements shift up the Sn segregation threshold, Pt has the highest impact. It has otherwise been shown that Co addition has a weak impact on morphological or electrical properties. Indeed, the surface morphology
CRediT authorship contribution statement Andrea Quintero: Investigation, Conceptualization, Methodology, Visualization, Writing - original draft. Patrice Gergaud: Methodology, Investigation. Jean-Michel Hartmann: Investigation, Writing review & editing. Vincent Reboud: Conceptualization, Supervision. Eric Cassan: Supervision. Philippe Rodriguez: Investigation, Conceptualization, Methodology, Visualization, Supervision, Writing - original draft. Acknowledgments The authors would like to thank Virginie Loup for her help in GeSn surface preparation and Nicolas Chevalier and Denis Mariolle for their support during the AFM measurements. This work was supported by the French National Research Agency (ANR) under the ‘‘Investissements d’avenir’’ programs: ANR 10-AIRT0005 (IRT NANOELEC) and ANR 10-EQPX-0030 (EQUIPEX FDSOI 11) and by the CEA DSM-DRT Phare project, France ‘‘Photonics’’. 5
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Impact of alloying elements (Co, Pt) on nickel stanogermanide formation Andrea Quintero a,b , Patrice Gergaud a , Jean-Michel Hartmann a , Vincent Reboud a , Eric Cassan b , Philippe Rodriguez a ,∗ a b
Univ. Grenoble Alpes, CEA, LETI, F-38000 Grenoble, France Centre de Nanosciences et de Nanotechnologies, CNRS, Univ. Paris-Sud, Université Paris-Saclay, C2N - Orsay, 91405 Cedex, France
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Keywords: Advanced metallization Ni NiCo NiPt Solid-state reaction Stanogermanide GeSn
ABSTRACT The impact of Pt or Co as alloying elements for Ni-based metallization of GeSn layers has been investigated. As far as the solid-state reaction is concerned, the overall phase sequence is the same for all metallizations: at low temperature, a Ni-rich phase is obtained; it is then consumed to form the low resistivity mono-stanogermanide phase. Nevertheless, the addition of an alloying element has an impact on Ni consumption, Ni-rich and monostanogermanide phases’ formation temperatures. Moreover, the addition of Co or Pt positively impacts Sn segregation by delaying this phenomenon. Co has a weak influence on morphological and electrical properties. On the other hand, Pt improves the surface morphology by delaying the Ni(GeSn) phase agglomeration and enhancing the process window in which the sheet resistance remains low.
1. Introduction Ge1-x Snx alloys have seen their interest growing in the past few years. These materials are envisioned for electronic applications such as source and drain stressors in Ge MOSFETs [1] and high mobility channels in pMOSFETs [2] or pTFETs [3]. Since the addition of enough Sn in the Ge lattice structure (about 10 at.%) results in a direct band-gap semiconductor, Ge1-x Snx layers have also been considered for optoelectronic applications such as p-i-n photodetectors [4,5], heterojunction LEDs [6,7] and, more recently, optically pumped lasers [8,9]. Whatever the targeted application, an efficient electrical ohmic contact is needed to obtain high performance devices. Ni-based contacts have been proposed and widely studied as they enable the formation of Ni(GeSn) intermetallics at relative low temperature with low values of contact (Rc ) and sheet (Rsh ) resistances [10, 11]. Nevertheless, Ni-based contacts exhibit a poor thermal stability linked to the precipitation of Sn atoms and the agglomeration of the Ni(GeSn) layer. These phenomena are enhanced by the increase of Sn content in the Ge1-x Snx layers [12,13]. To extend the thermal stability of Ni-based contacts, solutions have been proposed such as the cosputtering of Ni and Pt [14,15] or the addition of a Pt interlayer [16, 17]. As an alternative to Ni-based contact and in order to extend the thermal stability of contacts, Ti-based contacts have also been evaluated [18–20]. In this study, the impact of alloying elements, namely Co and Pt, on the Ni/GeSn system was investigated. The addition of Pt or Co should enhance the thermal stability of stanogermanides, given what is known
for the Ni/Si(Ge) systems [21–26]. Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 metallizations on Ge0.9 Sn0.1 layers were thus investigated. In particular, the impact of alloying elements on the solid-state reaction and the contact properties – in terms of surface morphology and electrical properties – was studied. The aim of this study was to highlight the opportunities of using alloying elements and CMOS-compatible metallization schemes to enhance the thermal stability of Ni-GeSn contacts. 2. Experimental GeSn layers with 10 at.% of Sn were epitaxially grown at 325 ◦ C, 100 Torr with Ge2 H6 and SnCl4 precursors on Ge-buffered Si (1 0 0) substrates in reduced pressure chemical vapor deposition (RPCVD) chambers [27]. The Sn composition was compatible with the optoelectronics applications described in the introduction. The thickness of the GeSn layers (60 nm) was low enough for them to be fully compressively strained on the Ge strain relaxed buffers (SRB) underneath. The layer thickness, 60 nm, and the Sn content, 10 at.%, were extracted from 𝜔–2𝜃 scans around the (0 0 4) XRD order. Surface preparation before metal deposition consisted in dips for 1 min in hydrofluoric acid (HF) diluted at 1% in water followed by rinses in deionized water and drying with a N2 gun – in order to remove particles and native oxides from the surface. Then, the deposition, at room temperature, of 10 nm of Ni, Ni0.9 Co0.1 or Ni0.9 Pt0.1 capped with 7 nm of TiN was performed by magnetron sputtering in a 300 mm cluster tool. To limit the regrowth of native oxides, the time interval between the surface preparation and
∗ Corresponding author. E-mail address: [email protected] (P. Rodriguez).
https://doi.org/10.1016/j.mssp.2019.104890 Received 11 September 2019; Received in revised form 18 November 2019; Accepted 16 December 2019 1369-8001/© 2019 Elsevier Ltd. All rights reserved.
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Fig. 2. In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni0.9 Co0.1 and Ge0.9 Sn0.1 .
Fig. 1. In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni and Ge0.9 Sn0.1 .
fluctuation of the hexagonal metastable 𝜖-Ni5 (GeSn)3 phase when the Ge(Sn) composition varied [28]. At 120 ◦ C, Ni was totally consumed. At 210 ◦ C, the diffraction peaks linked to the Ni-rich phase dissipate and new diffraction peaks appear at 42.4◦ , 45.3◦ and 53.4◦ . They correspond to the formation of the Ni mono-stanogermanide phase Ni(GeSn) with the following orientations: (2 1 0), (2 1 1) and (0 1 3). This phase remains stable until 675 ◦ C, a temperature at which the diffraction peaks vanish. This is due to a phase destabilization when reaching the NiGeSn–GeSn eutectic point. Even if the eutectic point between NiGe and Ge was reported to be at around 760 ◦ C [29], the addition of Sn in the Ge lattice might modify this temperature. In addition, as for all annealing equipment, the temperature indicated by the furnace and the absolute temperature might differ by a few tens of degrees. In order to study a potential impact of the addition of 10 at.% of Co on Ni-based metallization, the Ni0.9 Co0.1 /Ge0.9 Sn0.1 solid-state reaction was monitored by in-situ XRD. Patterns are shown in Fig. 2. The phase sequence is the same than for the Ni/GeSn system: the Ni0.9 Co0.1 (diffraction peak at 44.4◦ ) is consumed to the benefit of a NiCo-rich phase (NiCo)5 (GeSn)3 (diffraction peak at 46.5◦ ) which is in turn consumed to form the NiCo mono-stanogermanide phase (NiCo)(GeSn) (diffraction peak at 2𝜃 = 53.9◦ ). Nevertheless, a few discrepancies are seen. First, the NiCo-rich phase is present in films right after the deposition state. At high temperature, after the disappearance of the (NiCo)(GeSn) phase when reaching the NiGeSnGeSn eutectic point, another diffraction peak at 34.7◦ is observed. This diffraction peak was not detected during comparable in-situ XRD measurements on Ni/GeSn and NiCo/Ge systems. This peak is thus due to the concomitant presence of Co and Sn and the formation of CoSnx compounds. The impact of the addition of 10 at.% of Pt on the Ni-based solidstate reaction was also studied by in-situ XRD. Fig. 3 shows the contour map obtained for the Ni0.9 Pt0.1 /Ge0.9 Sn0.1 system. This system is by far the most complicated one among the three systems investigated in this study. A complete analysis of the Ni0.9 Pt0.1 / Ge0.9 Sn0.1 system is reported elsewhere [30]. At relative low temperatures, the phase sequence is similar to the one obtained with Ni or Ni0.9 Co0.1 : after Ni0.9 Pt0.1 consumption, the (NiPt)5 (GeSn)3 rich phase is formed. The later is consumed to form the NiPt monostanogermanide phase (NiPt)(GeSn). At around 350 ◦ C, a shift towards higher 2𝜃 angles is observed. This shift appears concomitantly with the appearance of new diffraction peaks at 39.2◦ and 42.0◦ . This is due to a rejection of Pt from the (NiPt)(GeSn) lattice and a transition towards Ni(GeSn) with the formation of PtSnx compounds with various stoichiometries and thermal stabilities [30].
the metallization of the samples was reduced to the minimum (less than 10 min). In addition, there was no air break between Ni(Co,Pt) and TiN capping layers’ deposition: samples were indeed kept at pressures less than 10−7 Torr during transfer between one chamber and another. The TiN layers preserved the Ni(Co,Pt) layers from atmospheric contamination. This way, solid-state reaction between Ni(Co,Pt) and Ge0.9 Sn0.1 was not impacted by oxygen or carbon contamination. The Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 phase formation sequence and crystalline evolution were followed by in-situ X-ray diffraction (XRD) using an Empyrean PANalytical X-ray diffractometer equipped with a copper (Cu K𝛼) source, a linear detector (PIXcel1D ) and a HTK 1200 Anton Paar furnace under secondary vacuum (pressure was maintained at about 10−6 bar to avoid atmospheric contamination). 𝜃–2𝜃 scans over a large range of diffraction angles (20-90◦ ) were performed with a 2◦ offset in 𝜔 in order to attenuate the symmetric (4 0 0) substrate reflection. Profiles were acquired from 50 up to 600 or 700 ◦ C, with 5 ◦ C steps and a scanning time of 7 min per temperature step. Phase analyses were performed with the HighScore Plus software from PANalytical, using the International Centre for Diffraction Data (ICDD) PDF-4 database. Furthermore, samples metallized with Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 using the same procedure as described above were ex-situ annealed for 30 s at temperatures ranging from 150 up to 550 ◦ C in a rapid thermal annealing (RTA) furnace with a N2 environment. Then, 𝜃–2𝜃 scans over a large angular range (20–90◦ ) using a Cu source were performed in a X’Pert Pro PANalytical X-ray diffractometer. The surface morphology and the electrical characteristics were assessed by tapping mode atomic force microscopy (AFM) with a Bruker FastScan equipment and sheet resistance (Rsh ) with a four-point probe meter. 3. Results & discussion 3.1. Study of the solid-state reaction The Ni/Ge0.9 Sn0.1 solid-state reaction followed by in-situ XRD is shown in Fig. 1 as a contour map. The 𝑦-axis shows the temperature and the 𝑥-axis, the diffraction angles. Peaks of TiN (1 1 1) at 36.4◦ , Ni (1 1 1) at 44.9◦ , Si (4 0 0) and Ge/GeSn (4 0 0) around 68◦ are observed. At low temperature (above 85 ◦ C), new diffraction peaks appear at 2𝜃 around 46.6◦ and 85.6◦ . These peaks are due to the formation of the 𝜖-Ni5 (GeSn)3 rich phase [28] with the following orientations: (1 1 0) and (3 0 0). The shift towards higher 2𝜃 angles observed at 130–140 ◦ C was linked to stoichiometry variations, i.e. lattice constant 2
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Fig. 5. Phase sequences of the reaction of 30 nm-thick Ni, Co or Pt films on Ge (0 0 1). Source: Adapted from [32].
Fig. 3. In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni0.9 Pt0.1 and Ge0.9 Sn0.1 .
soluble in Ni5 Ge3 or NiGe should decrease the nucleation barrier [31]. The phase sequences of Ni, Co and Pt germanides are depicted in Fig. 5. The NiGe phase can coexist with (i) CoGe, Co5 Ge7 and CoGe2 phases and (ii) Pt2 Ge, Pt3 Ge2 , PtGe, Pt2 Ge3 and PtGe2 phases [32]. NiGe being orthorhombic, all the orthorhombic phases might be soluble in NiGe. For the Pt/Ge system, all the intermetallic phases, except Pt2 Ge which is hexagonal, are orthorhombic. For the Co/Ge system, CoGe2 only is orthorhombic. Indeed, CoGe is monoclinic and Co5 Ge7 tetragonal. Thus, based only on solubility point of view, the fact that Co decreases the formation temperature of NiGe whereas Pt shifts it up towards higher temperatures cannot readily be explained. The Ni-rich phase, Ni5 Ge3 , coexists only with transition metals. The fact that the Co has an hexagonal structure, like Ni5 Ge3 while Pt has a cubic structure might explain why the addition of Co decreases the energy threshold for the Ni-rich phase. For the mono-germanide phase, if the impact of the alloying element on the nucleation barrier is not obvious, one should consider surface/interface effects. The Co–Ge, Ni–Ge and Pt–Ge bond strengths are equal to 230, 290 and 400–500 kJ mol−1 , respectively [33–35]. If the bonding between atoms is weaker (which is the case for Co), surface, interface and grain boundary diffusions are easier. As a consequence, atomic diffusion will be faster, resulting in a shift of the formation of intermetallics towards lower temperature with Co. The same reasoning would explain why the addition of Pt shifts the formation of intermetallics towards higher temperatures. In addition, it can be seen in Fig. 3 that, at around 120 ◦ C, the Ni0.9 Pt0.1 alloy peak shifts towards lower 2𝜃 angles, suggesting an increase of the alloy lattice parameter. The phenomenon is known as Pt snowplow and has already been discussed elsewhere [30]. This phenomenon might explain the delay noticed for the appearance of the Ni-rich phase with the addition of Pt.
Fig. 4. Impact of Co or Pt addition on Ni consumption, Ni-rich and monostanogermanide phases’ formation temperatures.
The study of Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 solid-state reactions has highlighted that, if the overall phase sequence is the same, discrepancies do exist. Fig. 4 summarizes the impact of Co or Pt addition on Ni consumption, Ni-rich and mono-stanogermanide phases’ formation temperatures. As far as Ni consumption is concerned, the addition of 10 at.% of Co does not impact the temperature at which the Ni is totally consumed. Indeed, Ni full consumption is observed at 120 ◦ C for both metallizations. On the other hand, the addition of 10 at.% of Pt pushes the Ni total consumption up to 140 ◦ C. Pt redistribution during thermal reaction, in particular the Pt snowplow effect, might explain this delay in total Ni consumption [30]. The formation of the Ni-rich phase is strongly impacted by the nature of the alloying element. For Ni metallization, the Ni5 (GeSn)3 phase appears at 85 ◦ C and is totally consumed at 210 ◦ C. When Co is added, the (NiCo)5 (GeSn)3 phase is observed since the as-deposited state. This phase is totally consumed at 210 ◦ C. On the other hand, the addition of Pt postpones the formation of the Ni-rich phase to 140 ◦ C and its total consumption to 240 ◦ C. Similar trends are observed for the temperature at which the mono-stanogermanide phase appears. Co addition tends to lower this temperature (200 ◦ C vs. 210 ◦ C for Ni) whereas Pt addition pushes this temperature up to 240 ◦ C. To sum up, the addition of Co lowers the formation temperatures of the Ni-rich and mono-stanogermanide phases whereas Pt addition delays the formation of these intermetallic phases. The impact of the alloying element’s nature on the formation temperature of intermetallics can be discussed according to various lines. From a solubility point of view, elements forming germanides that are
3.2. Impact on the Sn segregation There was, upon Co or Pt addition, a formation of CoSnx and PtSnx compounds at high temperatures (see the solid-state reactions described in Section 3.1). The formation of these kinds of compounds suggests the presence of Sn in the metallic state. However, no diffraction peak was linked to the Sn phase in Figs. 1 to 3. Indeed, beyond 230 ◦ C, tin is in a melted state and no diffraction peak can be detected with in-situ XRD analyses. Thus, XRD analyses were systematically performed on samples annealed by RTA to highlight the impact of the alloying element on the Sn segregation threshold. The Sn segregation threshold temperatures for each system, Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 metallizations on Ge0.9 Sn0.1 , are plotted in Fig. 6. As far as the Ni/Ge0.9 Sn0.1 system is concerned, no diffraction peak linked to the 𝛽-Sn phase was discernible at 300 ◦ C. Diffraction peaks associated with the 𝛽-Sn phase were however observed since 350 ◦ C. Thus, the Sn segregation threshold for this system is ranging between 300 and 350 ◦ C (hatched area in Fig. 6). In the same way, the 3
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Whatever the system studied, at low temperature (i.e. for T ≤ 300 ◦ C), samples exhibit cross hatch patterns along the ⟨1 1 0⟩ crystallographic direction. This pattern is due to the propagation of the 60◦ threading arms of misfit dislocations on (1 1 1) planes during Ge buffer and subsequent GeSn epitaxial growth [27,42]. This means that, in the as-deposited state, Ni-based and TiN layers are extremely smooth and their deposition does not modify the original surface morphology. This morphology is retained up to 300 ◦ C, meaning that the formation of the Ni-rich phase does not modify the original morphology. At 350 ◦ C, the surface morphology for Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems drastically changes, with a strong increase of the surface roughness. This change has been linked to the formation of the monostanogermanide phase and its subsequent agglomeration [30]. At higher temperatures, the roughness gradually increases due to the amplification of both phenomena: the mono-stanogermanide phase agglomeration and Sn segregation. The surface morphology evolution is very similar for Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems. Thus, the addition of Co as an alloying element does not really impact the surface morphology. The analysis of the Ni0.9 Pt0.1 /Ge0.9 Sn0.1 system highlights substantial differences in terms of surface morphologies. At 350 and 400 ◦ C, even if the morphology is changing with the appearance of cracks at the sub-micron scale on the surface, the cross-hatch is still visible. These cracks are linked to grain boundary grooving phenomenon at the early stage of agglomeration. The cross-hatch is no longer observable at 450 ◦ C and above, and the surface roughens. Fig. 7 clearly shows that the addition of Pt has a positive impact on surface morphology, by delaying the agglomeration of the Ni(GeSn) layer. An in-depth discussion can be found in a previous paper [30]. Briefly, the Pt–Ge bond energy being higher than the Ni–Ge one, Pt atoms will reduce the breakdown rates of chemical bonds which will reduce grooving and could explain the morphological stability improvement at high temperatures. Pt addition might also suppress or limit Ge out diffusion, which is beneficial for the morphological stability [35]. It has also been demonstrated that Pt addition reduces the grain size of the monostanogermanide phase [30], reducing agglomeration phenomena [43, 44].
Fig. 6. Sn segregation threshold extracted from XRD analyses of samples annealed by RTA.
Sn segregation threshold for the Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems are ranging between 350 and 400 ◦ C and 400 and 450 ◦ C, respectively. Thus, the addition of Co or Pt increases the Sn segregation temperature. The GeSn layers being grown by RP-CVD under kinetic conditions far from equilibrium, they are metastable. Indeed, GeSn is known for its poor thermal stability. Sn has been reported to segregate at the surface when the annealing temperature increases [36,37]. Sn segregation is linked to the formation of Sn clusters/precipitates in the GeSn layers and the subsequent migration of atoms towards the surface [36–38]. It can occur more easily near the surface because the surface free energy is higher than the bulk free energy [36]. Even if reducing the density of point defects limits the initial formation of Sn clusters/precipitates, Sn segregation is mainly governed by the migration/diffusion of atoms. Thus, the addition of alloying elements might impact (i) the diffusion of Sn atoms and (ii) the surface/interface energies involved in Sn segregation mechanisms. As reported in Section 3.1, the addition of Co or Pt leads to the formation of CoSnx and PtSnx compounds. On the other hand, the formation of NiSnx compounds was not observed. After the formation of Sn clusters/precipitates in GeSn layers, Sn atoms diffuse towards the surface. This diffusion can occur within the grains or, much more likely, at the grain boundaries of Ni(GeSn). In the presence of alloying elements, such as Co or Pt, the Sn diffusion might be slowed down by the formation of CoSnx and PtSnx compounds, resulting in a delayed Sn surface segregation. The second hypothesis concerns the modification of the surface/ interface energies by the alloying element. It is well known that the addition of alloying elements can impact the texture of silicides and germanides [39]. As seen in Figs. 1 to 3, the number of out-of-plane orientations for the mono-stanogermanide phase decreases with the addition of an alloying element: from 3 with pure Ni to only one orientation with the addition of Co and Pt. The formation of highly textured films has already been reported for Pt addition in Ni-based silicides [21,40]. They are expected to have a significantly lower interfacial energy compared to polycrystalline films. Thus, the formation of more textured mono-stanogermanide films with the addition of an alloying element can result in lower the surface/interface energies and have a positive impact on Sn segregation. Finally, Sn segregation has also been linked to strain at the interface [41]. The covalent radius of Ni (124 pm) being lower than the ones of Co (126 pm) and Pt (136 pm), the addition of these elements might modify the local strain and positively impacts the Sn segregation threshold. As seen in Fig. 6, the higher the element’s radius is, the higher the Sn segregation threshold is.
3.4. Impact on the electrical properties Fig. 8 shows the evolution of the sheet resistance after various RTA from 200 up to 550 ◦ C for the Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems. Whatever the system, from the as-deposited state and up to 300 ◦ C, sheet resistances (Rsh ) remain constant at around 28–35 Ω/sq. This temperature range corresponds to the Ni-rich phase domain. At 350 ◦ C, minimum values are reached for all systems: 23.1, 16.3 and 10.4 Ω/sq for the Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems, respectively. Differences are significant. The mono-stanogermanide layer obtained at 350 ◦ C with the addition of Pt as an alloying element is less resistive than the one obtained with Co and even less resistive than the one without any alloying element (pure Ni). These differences might be due to thickness variations, modifications of the intrinsic properties of the intermetallic layers by the alloying element or to the layer’s morphology itself. For Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems, a strong increase of the Rsh values is observed above 350 ◦ C. This behavior is directly linked to the agglomeration of the NiGeSn layer and to Sn segregation. Thus, as observed for the morphology, the addition of Co does not significantly impact the electrical properties. The decrease of the Rsh values at high temperature observed for the Ni0.9 Co0.1 /Ge0.9 Sn0.1 system might be linked to the formation of CoSnx compounds described in Section 3.1. On the other hand, the addition of Pt as an alloying element enhances the process window in which Rsh values remain low. Indeed, a plateau is observed between 350 and 400 ◦ C and the increase is moderate for 450 ◦ C. Beyond 450 ◦ C, the sheet resistance drastically increases, reaching values similar to that of the Ni/Ge0.9 Sn0.1 system.
3.3. Impact on the surface morphology Fig. 7 shows AFM images for Ni (top row), Ni0.9 Co0.1 (middle row) and Ni0.9 Pt0.1 (bottom row)/Ge0.9 Sn0.1 samples annealed at various temperatures. 4
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Fig. 7. Surface AFM images (5 μm x 5 μm scan sizes) of samples annealed at various temperatures. Ni (top row), Ni0.9 Co0.1 (middle row) and Ni0.9 Pt0.1 (bottom row)/Ge0.9 Sn0.1 systems. Scan directions along ⟨1 0 0⟩.
and the sheet resistance evolution as a function of annealing temperature are quite similar for Ni and Ni0.9 Co0.1 /Ge0.9 Sn0.1 systems. On the other hand, Pt positively impacts the surface morphology by delaying the Ni(GeSn) phase agglomeration and enhancing the process window in which the Rsh values remain low. Thus, even if both elements have a positive impact on Sn segregation, the impact on morphology, i.e. the ability to delay the agglomeration of the Ni(GeSn) phase, prevails in order to optimize the electrical properties. The addition of Co as an alloying element has a positive impact on Sn segregation, it also decreases the formation temperature of the Ni(GeSn) phase which might be valuable for some applications. Nevertheless, Pt addition, by enhancing the morphological and the electrical properties – mainly by delaying the Ni(GeSn) agglomeration and Sn segregation – is the most interesting option in terms of integration issues. It opens up the way towards a CMOS-compatible contact technology on GeSn. It was used on photonics components such as midinfrared photodiodes and LEDs. First results showed a strong reduction of the contact resistance compared to classical Ti/Al contacts and real improvements in terms of optoelectronic devices performances [45,46].
Fig. 8. Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for the Ni, Ni0.9 Co0.1 and Ni0.9 Pt0.1 /Ge0.9 Sn0.1 systems.
Declaration of competing interest 4. Conclusions The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
In this work, the impact of alloying elements, namely Co and Pt, on Ni stanogermanide formation and properties was investigated. As far as the solid-state reaction is concerned, the overall phase sequence is the same: at low temperature, a Ni-rich phase is obtained, the latter being consumed to form the mono-stanogermanide phase. Nevertheless, the addition of an alloying element has an impact on Ni consumption, Ni-rich and mono-stanogermanide phases’ formation temperatures. Based only on solubility point of view, the fact that Co decreases the temperature of formation of NiGe whereas Pt shifts it up towards higher temperatures cannot be readily explained. However, the impact of alloying element on atomic diffusion might affect surface/interface processes. The addition of an alloying element has also a positive impact on the Sn segregation threshold. Indeed, the addition of Co or Pt delays the Sn segregation temperature. The formation of CoSnx or PtSnx might slow Sn diffusion towards the surface and its subsequent segregation. Co and Pt addition might also lower the surface/interface energies, with a positive impact on Sn segregation. If both elements shift up the Sn segregation threshold, Pt has the highest impact. It has otherwise been shown that Co addition has a weak impact on morphological or electrical properties. Indeed, the surface morphology
CRediT authorship contribution statement Andrea Quintero: Investigation, Conceptualization, Methodology, Visualization, Writing - original draft. Patrice Gergaud: Methodology, Investigation. Jean-Michel Hartmann: Investigation, Writing review & editing. Vincent Reboud: Conceptualization, Supervision. Eric Cassan: Supervision. Philippe Rodriguez: Investigation, Conceptualization, Methodology, Visualization, Supervision, Writing - original draft. Acknowledgments The authors would like to thank Virginie Loup for her help in GeSn surface preparation and Nicolas Chevalier and Denis Mariolle for their support during the AFM measurements. This work was supported by the French National Research Agency (ANR) under the ‘‘Investissements d’avenir’’ programs: ANR 10-AIRT0005 (IRT NANOELEC) and ANR 10-EQPX-0030 (EQUIPEX FDSOI 11) and by the CEA DSM-DRT Phare project, France ‘‘Photonics’’. 5
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