Formation of epitaxial nanodots on Si substrates with controlled interfaces and their application

One ML of Si surfaces was oxidized under the condition: O₂ pressure of 10⁻⁵–10⁻⁴ Pa at 400–600 °C to form the ultrathin SiO₂ films. The films have amorphous morphology confirmed by STM and RHEED. Si or Ge atoms were deposited on the ultrathin SiO₂ films to form Si or Ge NDs. STM images of the Si and Ge NDs are shown in Figs. 1(a) and 1(b), where 8 ML Si and 5 ML Ge were deposited at 500 °C respectively. 5 nm spherical NDs seemed to be formed on Si substrates with an ultrahigh density of ∼2 × 10¹² cm⁻². RHEED patterns indicated that at deposition temperature higher than ∼450 °C, NDs were epitaxially grown on Si substrates, while at deposition temperature lower than ∼400 °C, RHEED of NDs exhibited Debye ring pattern revealing non-epitaxial growth of NDs.

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Formation mechanism was reported everywhere as follows.^26,27,38) At the first stage of Si or Ge deposition on the ultrathin SiO₂ films, deposited atoms diffuse on the surfaces until they react with the ultrathin SiO₂ films through the following reactions:

$$\begin{equation} \text{Si} + \text{SiO$_{2}$}\rightarrow \text{2SiO$\uparrow$}\quad \text{or}\quad \text{Ge} + \text{SiO$_{2}$}\rightarrow \text{SiO$\uparrow$} + \text{GeO$\uparrow$}. \end{equation} \tag{ 1 }$$

As a result, the nanovoids are formed in the ultrathin SiO₂ films. The distance between nanovoids is determined by diffusion length of deposited atoms until they react with the ultrathin SiO₂ films. In the case of Si and Ge, this value is approximately 10 nm resulting in the ultrahigh density (∼10¹² cm⁻²). Extra Si or Ge atoms are trapped at the nanovoids because nanovoids with dangling bonds are energetically stable sites for deposited atoms compared with the ultrathin SiO₂ film sites. In other words, the chemical potential of deposited atoms at the nanovoid sites is smaller than that on the ultrathin SiO₂ film. As a result, ultrasmall Si or Ge nuclei on nanovoids are formed, which are called as nanowindows. Then, nanowindows work as nucleation sites for ND growth by subsequent deposition resulting in the formation of ultrahigh density spherical NDs. When the deposition temperature is high ($\gtrsim$450 °C), nanowindows form sufficiently through reaction (1). As a result, the formed NDs contact with Si substrates through nanowindows leading to the epitaxial growth. On the contrary, when the deposition temperature is low ($\lesssim$400 °C), nanowindows do not penetrated into Si substrates due to insufficiency of reaction (1), resulting in non-epitaxial growth of NDs. As for epitaxial NDs, the limited nanocontact between NDs and substrates reduces the strain energy in NDs induced by lattice mismatch. The NDs are elastically-strain relaxed and have no misfit dislocation due to the small strain energy and spherical shape. This technique for ND growth can be applied to the ultrathin SiGeO_x films,³⁹⁾ as shown in Fig. 1(c).

We paid attention to the ultrathin SiO₂ film with Si or Ge nanowindows as a template for ND growth, and modified the aforementioned ultrathin SiO₂ film technique. As shown in Fig. 2, first, we predeposited a small amount of Si of Ge on the ultrathin SiO₂ film to form ultrahigh density nanowindows. Second, atoms A and B were deposited simultaneously on the ultrathin SiO₂ film with nanowiondws with flux ratio of 1 − x to x to form epitaxial A_1−xB_x NDs. Thus, the modified technique for ND growth can be applied to various materials. In this technique, we expected that the atoms A and B are trapped at the nanowidow sites, and then, ultrahigh density NDs were formed. Furthermore, NDs were epitaxailly grown on Si substrates by nancontact between NDs and substrates through nanowindows. In this case, as is the case with Si or Ge, limited contact between NDs and Si substrates enables elastic strain-relaxation in spherical NDs without introduction of misfit dislocations at their heterointerfaces.^32–34) In Fig. 3, STM images of the various kinds of NDs are shown.^30,31,40) The NDs are epitaxially grown on Si substrates, which were confirmed by RHEED patterns.

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In the case of β-FeSi₂ NDs on Si(001) substrates, the formation process is a little different from those on Si(111) substrates. 2 ML Si was deposited on the ultrathin SiO₂ films at 670 °C to form nanowindows. Fe and Si were codeposited on the ultrathin SiO₂ film at RT to form ultrahigh density amorphous iron silicide NDs. Then, the NDs were crystalized to epitaxial β-FeSi₂ ones by annealing at 450–650 °C. This technique is a solid phase epitaxy on the ultrathin SiO₂ films. ND shape and size strongly depended on the annealing temperatures.

The NDs exhibited the quantum confinement effects. The energy bandgaps of individual Ge and GeSn NDs were measured by scanning tunneling spectroscopy, which depended on ND size. The dependences are explained by the quantum confinement effect.

It was confirmed that the β-FeSi₂ NDs on Si(111) and (001) surfaces were indirect-transition semiconductor using STM-EFMS that is photoabsorption spectroscopy measurement with nanometer-sized special resolution⁴¹⁾ by comparing with bulk and film results.^42,43)

The rapid progress of information technology has largely relied on the development of nonvolatile memories. Recently, ReRAM has attracted much attention as one of the promising next generation nonvolatile memories.^51–55) It has advantages such as low power consumption, high switching speed. Operation of ReRAM is based on a resistive switching effect in various metal oxides. Fe oxide, composed of abundant and earth-friendly elements, shows a resistive switching effect.^56–58) We focused on Fe oxide grown on Si substrates because of their industrial advantages such as low cost, nontoxicity, and compatibility with existing Si process technology. Resistive switching in Fe oxide is reportedly caused by crystal structure changes between high resistance γ-Fe₂O₃ and low resistance Fe₃O₄ due to the oxygen vacancy movement. Compared with other materials with resistive switching effect, the resistance ratio between high and low resistance states, (OFF/ON ratio), was very poor (∼10).⁵⁷⁾ One of the causes can be the existence of crystal defects inducing the leakage current in low quality Fe oxide. Therefore, it is needed to form high quality Fe oxide epitaxially grown on Si substrates. We focused on single crystalline Fe oxide NDs epitaxially grown on Si substrates, where in a ND, there are no crystal boundaries and less crystal defects caused by crystal lattice mismatch.

There are many methods to form Fe oxide NDs, but in terms of epitaxial growth on Si substrates, there are not so many fabrication methods. One method can be the oxidation of the Fe NDs epitaxially grown on Si substrates. In addition to Fe oxide NDs for ReRAM, magnetic nanoparticles such as Fe and its oxide NDs have the potential to be applied as high-density magnetic storage media. Therefore, epitaxial growth of Fe and Fe oxide NDs on Si substrates is very meaningful for information technology. One of the difficulties in the epitaxial growth of Fe-based films/NDs on Si substrates is unintentional silicidation introducing the defective interfaces. For epitaxial growth, it is needed to use of buffer layers of insulators such as MgO or TiN^59–62) which are not suitable for ReRAM with flowing electric current. Up to now, there are no methods of direct epitaxial growth on Si substrates avoiding the silicidation.

The ultrathin SiO₂ film technique for epitaxial growth of iron silicide NDs on Si substrates has been developed where the interdiffusion between deposited Fe atoms and Si atoms in substrates was suppressed.⁴⁰⁾ However, Fe ND formation using ultrathin SiO₂ film technique has been not investigated minutely. Here, we present the formation technique of Fe-based NDs using ultrathin SiO₂ film technique.

We tried to develop the epitaxial growth technique of Fe NDs using the ultrathin SiO₂ films with Si nanowindows. 1 ML Si was deposited on the ultrathin SiO₂ films on Si(111) substrates at 600 °C to form the Si nanowindows. Fe (12–36 ML) was deposited on the ultrathin SiO₂ films with Si nanowindows at ∼100–350 °C. The STM images in Fig. 6 show the NDs grown under the various Fe deposition conditions in the case of the nanowindow formation in the ultrathin SiO₂ films at 600 °C. The average ND diameter and the density of formed NDs were summarized in Figs. 6(g) and 6(h). The ND diameter was ranging from 3 to 7 nm, and ND diameter became larger with increase of Fe deposition amount. The ND density is ultrahigh [(2–5) × 10¹² cm⁻²].

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Figures 7(a)–7(f) show the RHEED patterns of the NDs on Si(111) substrates grown under various Fe deposition conditions in this nanowindow formation case (Si deposition at 600 °C). Electron incident direction in RHEED observation was 〈112〉_Si. In the case of NDs formed by deposition of 12 ML Fe at ∼100 °C, RHEED pattern was too obscure due to the poor crystallinity and the small volume, while in the case of other NDs, diffraction spots were observed indicating that ultrahigh density NDs shown in Figs. 6(a)–6(f) were epitaxially grown on Si substrates. Closer investigation of RHEED patterns enabled the clarification of the crystal structure of formed NDs. For ε-FeSi NDs epitaxailly grown on Si substrates with ε-FeSi(111)/Si(111), the diffraction spots are appeared at the 1/3 position of the 1 × 1 Si diffraction streaks in RHEED patterns. The RHEED patterns of Fe₃Si, c-FeSi, and Fe NDs, which are similar to each other, are shown in Fig. 7(g). If NDs are composed of pure Fe ones epitaxially grown with Fe(111)/Si(111) without iron silicides, the pattern of bcc Fe shown in schematic in Fig. 7(g) appears, where the 111 for c-FeSi or 222 for Fe₃Si spot pointed by an arrow are not observed. For the NDs formed by Fe deposition at 350 °C, the 1/3 spots pointed by arrows were appeared in RHEED pattern, which indicated that these NDs included ε-FeSi. In the case of at lower Fe deposition temperature (100–300 °C), 1/3 spots did not appear, but it is not clear whether the silicide spots (111 for c-FeSi or 222 for Fe₃Si) existed which are shown by an arrow in Fig. 7(g). Therefore, we investigated the line profiles between points A and B in RHEED patterns in Figs. 7(c), 7(d), and 7(f). The line profiles are show in Fig. 7(h). The peaks were observed as pointed out by dotted circles in Fig. 7(h), which corresponded to 111 for c-FeSi or 222 for Fe₃Si indicating the existence of Fe silicide. By this means, in the case of nanowindow formation by Si deposition at 600 °C, the formation diagram was acquired as shown in Fig. 8. The epitaxial NDs were not composed of pure Fe, but included Fe silicide. This result indicated that in the case of nanowindow formation by Si deposition at 600 °C, the interdiffusion of Fe and Si in substrates took place. It was considered that interdifussion occurred through the nanowindows in ultrathin SiO₂ films at ∼100 °C not directly beyond the ultrathin SiO₂ films according to the report that Fe atoms were not able to diffuse beyond the ultrathin SiO₂ films to Si substrates at low temperature.⁶³⁾

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To prevent the interdiffusion of Fe and Si, we tried to reduce the nanowindow size by decreasing Si deposition temperature for forming nanowindows in the ultrathin SiO₂ films. Nanowindows were formed through the reaction shown by Eq. (1). Therefore, at lower Si deposition temperature, the reaction speed is reduced, resulting in smaller nanowindows. We formed nanowindows in the ultrathin SiO₂ films by 1 ML Si deposition at low temperature of 450 °C, and then deposited Fe deposition (12–36 ML) at ∼100 °C to form NDs. Figures 9(a)–9(c) show the RHEED patterns of the NDs indicating the diffraction pattern of epitaxially grown bcc-Fe. To clarify whether Fe silicide is included, we investigated the line profile between points A and B in RHEED [Fig. 9(c)], which is shown in Fig. 9(d). There is no peak corresponding to 111 spot for FeSi or 222 spot for Fe₃Si. This result indicates that the formed NDs were composed of Fe within the RHEED observation. Figures 9(e)–9(g) show the RHEED patterns of the NDs in the case of nanowindow formation temperature of 400 °C. Debye rings appeared indicating non-epitaxial NDs.

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As mentioned above, we aimed at changing nanowindow size by tuning Si deposition temperature for nanowindow formation. As a result, the strong dependence of ND growth on the Si deposition temperature was found. In the case of high nanowindow formation temperature (600 °C), Fe silicide such as Fe₃Si and c-FeSi were formed even when the Fe was deposited at low temperature (∼100 °C). In this case, the nanowindows were expected to be large to cause the interdiffusion of Fe and Si through them, as shown in Fig. 10. To make the nanowindows smaller, we reduced the nanowindow formation temperatures from 600 to 450 °C. In this case, the silicide were not observed indicating that the interdiffusion did not occur through nanowindows. As a result, pure Fe NDs were epitaxially grown on Si substrates. The epitaxial growth means that NDs contacted with Si substrates through nanowindows. On the other hand, at lower nanowindow formation temperature of 400 °C, the non-epitaxial NDs were formed. This is because nanowindow formation through reaction (1) did not proceed sufficiently at the lower temperature. As a result, the contact between NDs and Si substrates was not enough leading to non-epitaxial growth. The appropriate size of nanowindow is necessary for pure Fe NDs to grow epitaxially on Si substrates. In this case, the nanowindows formed by Si deposition at 450 °C were expected to have the most appropriate condition.

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36 ML Fe NDs were epitaxially grown on Si substrates using the ultrathin SiO₂ films with nanowindows formed by 1 ML Si deposition at 450 °C. The NDs were exposed to O₂ at 1 atm and RT. From STM image of the oxygen-exposed NDs, the ND density was found to be maintained at 2 × 10¹² cm⁻², while RHEED patterns showed changes from bcc crystal to an amorphous structure after exposure to oxygen. To estimate O₂ composition in the oxygen-exposed NDs, ex-situ X-ray photoelectron spectroscopy (XPS) were measured, revealing that oxygen-exposed NDs were oxidized at RT to be Fe₂O₃ NDs. XPS also revealed the existence of SiO₂.

Figure 11(a) shows cross sectional transmission electron microscope (HRTEM) image of the oxygen-exposed Fe NDs. HRTEM images showed the existence of amorphous NDs as indicated by the dotted circles in Fig. 11(a). The results of XPS and RHEED suggested these NDs were amorphous Fe₂O₃. In the layer beneath the Fe₂O₃ NDs, a second amorphous layer was observed to be SiO₂. The layer below the SiO₂ films was observed to be composed of crystalline and amorphous regions. FFT analysis of HRTEM images revealed that the crystal layer was Fe₃Si epitaxially grown on the Si substrates. The amorphous layer below the crystalline Fe₃Si appeared to be amorphous Fe silicide (a-silicide) resulting from incomplete silicidation.³⁵⁾

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When ultrathin SiO₂ films without Fe NDs are exposed to the atmosphere at room temperature, ultrathin SiO₂ films protects the further oxidation of Si substrates.^64,65) Additionally, before oxygen exposure, Fe deposition at ∼100 °C did not lead to silicidation on the ultrathin SiO₂ films with Si nanowindows formed at 450 °C as mentioned in the previous section. This indicated that oxygen exposure of the Fe NDs caused the Si oxidation and alloying of the Si and Fe at room temperature, as shown in Fig. 11(b). At the first stage of oxygen exposure, oxygen atoms were considered to go into Fe NDs and Si can be oxidized slightly near the interface with Si substrates. In the case of Fe films on Si that include a small proportion of oxygen (<2%),⁶⁶⁾ there is a report that Fe atoms were able to diffuse into Si substrates owing to the presence of small amounts of oxygen atoms. If this mechanism⁶⁶⁾ can be applied to our case, Fe atoms in Fe NDs including oxygen atoms diffuse into Si substrates resulting in the silicidation. In this framework, the aforementioned Fe diffusion accompanied by SiO₂ formation and silicidation should occur predominantly through nanowindows just below NDs, where NDs and Si were directly contacted. Furthermore, the existence of the a-silicide layer below the crystalline silicide layer can be reasonably explained by this mechanism, considering this amorphous layer as the reaction surface of silicidation.

To prove this mechanism, we tried to prevent Fe diffusion into Si in the presence of oxygen. We used Ge nanowindows, namely Ge nulcei on nanovoids, in the ultrathin SiO₂ films instead of Si nanowindows by depositing 5 ML Ge onto the ultrathin SiO₂ film at 620 °C. Fe was deposited on the ultrathin SiO₂ films with Ge nanowindows at 100 °C. STM and RHEED observations demonstrated that bcc-Fe NDs of several nanometer in diameter were epitaxial grown on Si substrates. The bcc-Fe NDs were exposed to oxygen at 1 atm, and then, the NDs of ∼10 nm in diameter were observed. A HRTEM image of oxygen-exposed Fe NDs displayed that flattened Fe oxide NDs were formed on epitaxial Ge nuclei on Si substrates. In the area of that Ge nuclei existed, there was no SiO₂ layer and no silicidation layer below the Ge nuclei. In areas without Ge nuclei, the 2-nm-thick SiO₂ above flat layers of 1 nm silicide were observed, but the thicknesses of silicide layer was smaller than the Si nanowindows case. This indicated that when Ge nuclei were present between the Fe NDs and Si substrates, oxygen-assisted Fe diffusion into Si did not occur, and silicidation and SiO₂ formation were suppressed.³⁵⁾

This demonstrated that the nanometer-sized interface between NDs and the substrates had remarkable influence on the interfacial reaction between Fe NDs and the substrate. Furthermore, control over the nanometer-sized interface is important for the use of Fe-based NDs on substrates with sharp interfaces, which could be an important general principle in system of NDs on substrates.

As mentioned in the previous section, Ge nanowindows work as barrier for silicidation between Fe and Si in substrates. In addition to the interest in the aforementioned Fe/Ge system, iron germanide has drawn much attention. Recent reports have invoked observations of Skyrmions in ε-FeGe(Si), that is a cubic B20 structure with chirality.^23–25) Another crystal structure of iron germanide, Fe_1.7Ge with a hexagonal B8₂ structure, has also been studied as a ferromagnetic material.⁶⁷⁾ In addition to the films, nanostructures of iron germanide on Si substrates are of great interest because nanostructuring can yield new properties and new physics. However, epitaxial growth of Fe(Ge)/Si with sharp interfaces is difficult because the interfaces are easily disordered due to silicidation (germanidation),^68–71) leading to non-controllable electronic states and a dead layer at the interface. Furthermore, the crystal structure control in this system is difficult owing to the complicated phase diagram. Therefore, a nanostructure formation technique with a well-controlled interface and crystal structure is crucially needed for this system.

We have developed an epitaxial growth technique for Fe NDs on small Ge nuclei³⁵⁾ by applying an ultrathin SiO₂ film technique. Now, we focus on use of these Ge nuclei as germanidation sites for epitaxial growth of iron gemanide NDs on Si substrates with sharp interfaces. We find that the crystal structure can be selected by controlling status of Ge nuclei.

We formed two kinds of Ge nanowiondws (Ge nuclei on nanovoids) in the ultrathin SiO₂ films by Ge deposition on the ultrathin SiO₂ films at 500 and 630 °C. Ge deposition at higher temperature get nanovoid size larger. Fattened Ge nuclei on large nanovoids were formed at Ge deposition temperature of 630 °C because of the larger naowindows at high deposition temperature, while at the Ge deposition temperature less than 600 °C, spherical Ge nuclei were formed. It was confirmed by RHEED and TEM that the flattened Ge nuclei were strained (s-Ge) due to the large contact area with substrates, while spherical Ge nuclei were elastically-strain-relaxed without misfit dislocation (r-Ge) owing to the small contact and its spherical shape. The lattice constant of s-Ge (r-Ge) is close to bulk Si (Ge), which is confirmed by RHEED pattern. This is illustrated in Fig. 12(a). To form iron germanide NDs, we deposited 5 and 7 ML Fe on the two kinds of Ge nuclei (s- or r-Ge nuclei) at ∼280 °C, which is a relatively high temperature compared with Fe ND case (∼100 °C). STM images demonstrate that the formed Fe-based NDs were flattened and spherical in the case of the s- and r-Ge nuclei, respectively. The ND density was the same as that of the Ge nuclei in both cases indicating that the Ge nuclei act as nucleation sites for Fe-based ND formation. The RHEED patterns implied that Fe_1.7Ge with the B8₂ structure⁶⁷⁾ was epitaxially grown on the Si substrates for r-Ge NDs, while ε-FeGe with the B20 structure was epitaxially grown on the Si substrates for s-Ge NDs. Fe_1.7Ge with the hexagonal phase is thermodynamically stable for a narrow stoichiometry interval ranging approximately from Fe₂Ge to Fe_1.5Ge. This is illustrated in Fig. 12(b). This can be explained by considering that the elastic strain energy largely contributed to free energy variation between various crystal structures of Fe-based NDs because the lattice mismatches in epitaxial growth on Si(111) between ε-FeGe and Si and between Fe_1.7Ge and Ge are both almost zero. This means that the crystal structure of iron germanide is determined by the elastic strain energy during the first stage of epitaxial growth, namely nucleation. As a result, by controlling the strain state, the lattice constant of the Ge nuclei working nucleation sites, the crystal structure of the epitaxially grown iron germanide could be selected.

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