Sn Spheres Embedded in a SiO2 Matrix: Synthesis and Potential

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Sn spheres embedded in a SiO matrix: synthesis and potential application as self-destructing materials Vu Xuan Hien, and Young-Woo Heo ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b05961 • Publication Date (Web): 27 Jul 2016 Downloaded from http://pubs.acs.org on August 8, 2016

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ACS Applied Materials & Interfaces

Sn spheres embedded in a SiO2 matrix: synthesis and potential application as self-destructing materials Vu X. Hien1, Young-Woo Heo2* 1

School of Engineering Physics, Hanoi University of Science and Technology, 1 Dai Co Viet

Street, Hanoi, Vietnam 2

School of Materials Science and Engineering, Kyungpook National University, Daegu 702-701,

Republic of Korea KEYWORDS silicon dioxide, tin, thermal evaporation, self-destructing materials, nanoexplosive material

ABSTRACT. We introduce a simple process for the fabrication of SiO2 films embedded with βSn-rich nano/microspheres. Sn spheres with maximum and minimum sizes of 10 µm (near the SiO2 surface) and 5 nm (at the Si/SiO2 interface) were grown within a 0.7–5.7 µm-thick SiO2 layer by evaporating SnO powders onto an Si (100) substrate for 1–600 min at 600–900 oC and 0.001–5.0 Torr. A possible growth mechanism of these materials is discussed. The currentvoltage characteristics of the as-fabricated samples were investigated in order to identify potential applications. During these tests, small flashes of light and the presence of damaged areas were observed at the oxide surfaces of the samples using an optical camera and a field

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emission scanning electron microscope, respectively. The electrical breakdown and shutdown of the devices observed in the current-voltage curves were attributed to the destruction of the SiO2 surface. In addition, the current-time responses show that the size of the damaged regions can be controlled by the voltage and duration of the applied stress, and are independent of the size and shape of the electrodes. The present materials thus possess great potential for applications in selfdestructing devices.

INTRODUCTION Nanoexplosive methods represent an interesting branch of nanoscience and nanotechnology, with a great potential in various applications. They were established in 1992, when McCord and coworkers noticed a “flash of light with an audible pop” during the reaction between concentrated HNO3 and porous silicon.1 A significant breakthrough in the field of nanoexplosive materials was achieved in 2002, when Mikulec et al. reported a solid-state explosive using porous silicon and a nitrate solution.2 The number of studies in this area has been rapidly growing since then, and many possible applications of nanoexplosive materials have been proposed, from the analysis of atomic emission spectroscopy data in portable devices, to the propulsion of miniature chemical devices (mini rockets), to the protection of the design of specialized chips (self-destructing chips), to smart ignition systems for conventional explosives, and propulsion systems for micro-electromechanical systems.3–8 The main drawback of using porous silicon as a nanoexplosive material is the oxidation of Si upon exposure to ambient air: the formation of SiO2 on the porous silicon substrate can hinder or inhibit the explosion.2,4 The fabrication of Sn nano/microcrystals in a SiO2 matrix has been studied for more than a decade, with the purpose of developing novel devices.9 Many studies have focused on the properties of Sn-implanted SiO2 structures, mainly using photoluminescence (PL).10–15 The


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present study introduces a simple synthesis route, based on an evaporation method, for the fabrication of β-Sn submicron spheres embedded in a SiO2 film. We show that the thickness of the SiO2 layer and the dimensions of the Sn submicron spheres can be tailored by tuning the working pressure, treatment time, and temperature during the evaporation. Finally, we demonstrate the self-destructing behavior of a prototype device using the as-synthesized material.

RESULTS AND DISCUSSION Morphology and structure of SiO2 film embedded with β-Sn submicron spheres Surface and cross-sectional FE-SEM images of the silicon substrate after treatment for different times (1–600 min) and at different pressures (0.001–760 Torr) and temperatures (600– 900 oC) are shown in Fig. 1. At atmospheric pressure, the surface was covered by numerous nanoparticles with an average dimension of 20 nm (Fig. 1(a)). The size of the particles significantly increased with decreasing pressure, and they changed into irregular islands (average size of 400 nm) at 50 Torr (Fig. 1(b)). Interestingly, the microspheres were well formulated on the sample surface at a working pressure equal to or lower than 5 Torr (Fig. 1(c-f)). At the higher vacuum level of 0.001 Torr, the average sphere size increased to approximately 1.2 µm, and the spheres were distinctly scattered on the sample surface. Surface and cross-sectional images of samples treated at 700 oC and 0.1 Torr for growth times of 1–600 min are presented in Fig. 1(gj). The figures clearly indicate that the thickness of the film (between 720 and 3550 nm) increases with the treatment time. A gradient in the sphere size is observed going from the interface between Si substrate and as-grown film to the film surface, with the largest spheres found near the film surface. Moreover, no hemispherical grains are visible on the film surface.


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The growth seems to saturate after 300 min of treatment, because the film surface (exhibiting many cracks) and the thickness of the samples treated for 300 and 600 min are quite similar. The influence of the temperature on the growth process of the samples (with a working pressure of 0.1 Torr and a treatment time of 60 min) is illustrated in Fig. 1(k-n). At 600 oC, a 300 nm thin film embedded with spherical grains is grown on the Si substrate. The film thickness increases to 1710 nm and approximately 5140 nm at 700 oC and 800 oC, respectively. Several grains presented on the film surface and many microellipsoids inside the film are observed in the sample treated at 800 oC (Fig. 1(m)). In the case of the sample grown at 900 oC, the film exhibits a wavy surface containing many islands of undefined shape, and large (~ 7–10 µm) spherical grains are also found above the surface. The elemental mapping images in Fig. 2 illustrate the atomic distributions of O, Si, and Sn in the sample grown at 700 oC and 1 Torr for 60 min. The thick film is composed of Si and O, while the spherical grains comprise a large amount of Sn. The energy dispersive X-ray (EDX) spectrum and the elemental line-scan analysis images of the sample in Fig. S3 confirm that the thick film and the spherical grains consist of SiO2 and Sn, respectively. The presence of metallic Sn can be understood on the basis of the following SnO disproportion reactions:16 3SnO ( s ) → Sn2O3 ( s ) + Sn(l ) (fast reaction)

(1)

2 Sn2O3 ( s ) → 3SnO2 ( s ) + Sn(l ) (slow reaction)

(2)

Moreno et al. reported that this two-step reaction can take place above 300 oC, with the fastest rate achieved in the temperature range of 400–500 oC.17 The same study also found that the breakdown ratio of SnO2 seems to be accelerated as the oxygen partial pressure is reduced. Although the growth process was carried out in N2 atmosphere, the film containing Sn


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nano/microspheres consists of SiO2 instead of SnO2. This is a possible consequence of the decomposition of SnO2 in reduced atmosphere to form O2(g), which supports the growth of SnO2 nano-whiskers and nano-wires.18,19 The decomposition of SnO2 may result from the following reactions occurring at defect sites:20 4+ SnO2 ↔ SnSn + OO2 − + VO2 − + 1 O2 2

(3)

SnO2 ↔ Sni + O2 ( g )

(4),

where VO and Sni denote an oxygen vacancy and an interstitial tin atom, respectively. The subsequent oxidation of Si then results in the formation of the SiO2 layer. It should be noted that the growth rate of this oxide layer is much faster than that reported in previous studies on the oxidation of Si in dry/wet oxygen.21,22 This discrepancy may be attributed to the large amount of Sn nano/micro-spheres present, which induces very large stress on the SiO2 layer during its growth. Therefore, the actual as-grown SiO2 layer is similar to a thick beehive-like structure. The transmission electron microscopy (TEM) data for the submicron spheres in Fig. 3(a) were assigned to metallic β-Sn with a tetragonal structure (International Centre of Diffraction Data (ICCD) powder diffraction file (PDF) no. 040673). The corresponding fast Fourier transform (FFT) pattern in the inset of Fig. 3(a) supports an interplanar spacing of 0.291 nm, matching the calculated spacing of the (200) plane of tetragonal Sn (a = b = 5.831 Å and c = 3.182 Å). In addition, numerous β-Sn nanocrystals with diameter in the range of 5–15 nm were found at the interface between SiO2 layer and Si substrate (Fig. 3(b)). Fig. 3(b) also shows that the β-Sn nanocrystals were embedded in the Si matrix without introducing any distortion in the Si lattice. The corresponding Miller indices and interplanar spacings of Si (220) and Sn (200) planes are shown in the FFT patterns of Fig. 3(c)-3(e). The structure of β-Sn in Si matrix in this study is


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fully consistent with previous work by Lei et al. and Lopes et al.23,24 Moreover, a simple explanation for the formation of β-Sn nanocrystals in the Si matrix was proposed24 using the SnSi equilibrium phase diagram.25 Based on this diagram, the concentration of Sn in the solid solution of Sn and Si increases with the temperature, e.g., ~ 0.03 at.% at 600 oC, up to a maximum of ~ 0.1 at.% at 1065 oC. Because the as-grown material is composed of β-Sn and SiO2, each element can be easily removed by stripping in solution, as shown in Fig. 4. Interestingly, after removing the β-Sn/SiO2 layer the Si surface exhibits numerous nanoscale pores with pyramidal shape, as shown in Fig. 4(e) and 4(f). According to the field emission scanning electron microscopy (FE-SEM) and scanning probe microscopy (SPM) data in Fig. 5, the quantity as well as the size of the pyramidal pores increased at higher treatment temperatures. Pores with a maximum size of 25 nm are partially visible on the Si surface of the sample treated at 600 oC, whereas the pore sizes of the samples treated at 700 and 800 oC are 50 and 180 nm, respectively. In addition, the penetration angle of Sn in the Si matrix is approximately 25.2o at 600 oC, indicating {311} planes intersecting the (100) surface. Therefore, the pyramidal pore shapes mentioned above for this sample are {311} facets. This result is similar to the reported gross surface rearrangement observed when 2.5 molecular Sn layer are deposited and annealed (at roughly 600 oC) on a Si (100) substrate, which also exposed {311} facets.26 For samples treated at 700 and 800 oC, the penetration angle is almost 54.7o, corresponding to the {111} pyramidal facets.

Growth mechanism of the SiO2 film embedded with β-Sn spheres The gradual formation of the SiO2 film embedded with Sn nano/microspheres is illustrated in Fig. 6. Depending on the vacuum level, the initial stage may involve the evaporation of Sn


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derived from the disproportion reaction described by Equations 1 and 2. This may result in the formations of liquid Sn nuclei on the Si substrate. Nano/microscale Sn islands can be formed during the deposition of Sn vapor, upon coalescence of the Sn nuclei (Fig. 6(a)). In the following, these islands will be labeled “surface tin” or “Snsurface” islands, which correspond to the hemispheres shown in Fig. 1(h). Fig. 1(e) and 1(f) also reveal that Snsurface islands are partially formed on the surface of the sample grown at p = 0.1 Torr, whereas the surface of the sample treated at p = 0.001 Torr is fully covered with Snsurface islands of larger size. This implies that the evaporation of liquid Sn can take place either before the oxidation of Si at high vacuum levels, or in the initial stage of the formation process. At a lower vacuum level ( p ≥ 0.1 Torr), the evaporation of Sn can only proceed at elevated temperature ( T ≥ 700 oC). In this case, the oxidation of the Si surface by oxygen decomposed from SnO2 (Equation 3 and 4) at around 600 o

C should be considered as the initial process. Therefore, no Snsurface islands are found in the

sample fabricated at p > 0.1 Torr (Fig. 1(d)). Depending on the treatment temperature, the oxidation process of Si may follow two modes. As shown by Enta et al., below 650 oC, the oxidation follows a random pattern without the formation of two-dimensional (2D) islands, whereas above 650 oC, it proceeds via 2D nucleation.27,28 This may explain the formation of the SiO2 particles and islands in the samples grown at 700 oC with p = 760 and p = 50 Torr (Fig. 1(a) and 1(b), respectively). Based on the evidence that the oxidation proceeds through the inward movement of the oxidant species rather than the outward movement of silicon, Deal and Grove suggested the following three-step process:21 (i) an oxygen species is transported to the outer surface of Si where it reacts or is adsorbed; (ii) the species is then transported across the oxide film towards the silicon substrate; (iii) it reacts at the silicon surface to form a new layer of


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SiO2. Although this process is of fundamental importance, it is not sufficient to explain the formation of the thick oxide layer in our sample. The evaporation and diffusion of Sn into the SiO2 layer during oxidation may represent the main cause of the increased SiO2 thickness in the present case. The substitutional-interstitial diffusion mechanism29 can be used to describe the diffusion of Sn to SiO2 or Si structures as follows: incident Sn atoms join SiO2 or Si structures as impurities; under elevated temperature, the impurity atoms become interstitials and move rapidly through the structure until they find voids or vacant lattice sites. Subsequently, they can rejoin the lattice and becomes substitutional atoms. SiO2/Si can be decomposed to form volatile SiO according to the following reaction Si(s) + SiO2(s) → 2SiO(g) at a temperature higher than 600 oC and low oxygen partial pressure.30,31 Therefore, silicon oxide decomposition can lead to the formation of microvoids in the oxide film, which can be easily occupied by Sn atoms.32 The diffusion of Sn atoms is restricted to the interface between SiO2 film and Si substrate. The interface may become a secondary ground for the nucleation of new Sn islands, as illustrated in Fig. 6(b). During the diffusion of Sn atoms, Sn nuclei or islands can be grown to form Sn submicron droplets (Sninternal) inside the SiO2 film (Fig. 6(c)) via Ostwald ripening.33 A study of the diffusion profile associated with the substitutional-interstitial diffusion mechanism carried out by Zahari et al.29 may explain why large and small Sn droplets are distributed near the oxide surface and the SiO2/Si interface, respectively. An alternative explanation is based on the well-known flexibility of the Si-O-Si angle (109.5o), which makes the Si-O-Si bonds easy to bend, stretch or rotate.34 These bonds can be influenced or broken by physical stress in the film, such as the interfacial stress generated by the volume expansion from Si to SiO2.35 Therefore, the large stress caused by mass diffusion of Sn into the SiO2 film can


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result in the breakage of Si-O-Si bonds and a large lattice distortion, which in turn may allow Sn and O2 to move easily across the oxide film. On the other hand, the diffusion of Sn and O2 can increase the oxide film density, forcing the oxide to release stress either by transporting Sn atoms to new layers of SiO2 generated by oxidation at the SiO2/Si interface, or by stretching the oxide layer near the film surface, because higher stress tends to be associated with a denser oxide film.36 When the oxide film is thick enough, the stress is preferentially released at the oxide surface, due to the ineffective oxidation of Si at the SiO2/Si interface. This may result in a density of the oxide layers close to the interface higher than that of regions near the film surface. In turn, this may determine a lower diffusion rate of Sn and O2 atoms from the surface to the SiO2/Si interface, and result in the Ostwald ripening of Sn occurring more easily near the oxide surface rather than in the region close the SiO2/Si interface. These effects explain why the larger Sn microspheres are observed near the oxide surface, whereas small Sn nanoparticles are found in the region close to the SiO2/Si interface (cross-sectional images in Fig. 1 and TEM image in Fig. 3). Moreover, the distortion and expansion of the SiO2 layers near the film surface may not only lead to the SiO2 layers overlapping with the Snsurface islands, but also generate cracks on the film surface, as in the samples grown at 700 oC and 0.1 Torr for 5 and 10 h (Fig. 1(i) and 1 (j), respectively). The processes inside the oxide film probably stop when the oxygen can no longer reach the SiO2/Si interface, so that an equilibrium is reached between the stress generation and relaxation, in which the diffusion process is prohibited. In situation, the growth can continue above the oxide surface, where Snsurface islands can become growing sites to form Sn grains (Fig. 6(d)). It is known that the density of β-Sn (7.29 g/cm3) is higher than that of amorphous SiO2 (2.196 g/cm3).37 Therefore, gravity may be a possible factor that leads the distribution of the


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microspheres inside the SiO2 layer. Under the gravity force, the β-Sn liquid may be pulled toward the SiO2 surface (the Si substrate was placed above the boat facing the SnO powder) and subsequently results in the formations of big Sn microspheres near the SiO2 surface and small Sn nanoparticles near the Si/SiO2 interface. This explanation is reasonable only if the Sn liquid can freely move inside the SiO2 layer. Nevertheless, the displacement of the Sn droplet is quite limited because the amorphous SiO2 still remains a soft solid layer at 900 oC. Therefore, the gravity can only result in the distortion of the Sn droplet. Indeed, the observation of many microellipsoids (with the longest diameter is above 2 µm) instead of microspheres in the sample treated at 800 oC (Fig. 1(m)) is a strong evidence for this hypothesis.

Electrical properties of SiO2 film embedded with β-Sn spheres In order to identify possible applications of these systems, the electrical properties of samples with an oxide layer thickness of 720 nm (sample A treated at 700 oC and 10-1 Torr for 1 min), 1710 nm (sample B treated at 700 oC and 10-1 Torr for 60 min) and 1969 nm (sample C treated at 700 oC and 10-3 Torr for 60 min) were investigated. Current-voltage (I-V) measurements of the samples were conducted at a positive voltage polarity (Fig. 7(a-c)). Five measurements were made for each sample, with an interval of one minute between subsequent measurements. The first I-V curve of sample A (Fig. 7(a)) shows an electrical breakdown, as indicated by the two steps at 8 and 16 V. The remaining I-V measurements of this sample denote conductive characteristics. In the first I-V curve of sample B (Fig. 7(b)), a current shutdown is observed at a voltage of approximately 14 V, after which the current is dropped from approximately 10-2 A down to roughly 10-5 A with the voltage increase from 14 to 20 V. The breakdown and shutdown voltages which characterize the I-V signal of both samples suggest that the samples were


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damaged during the first measurement. Recorded videos of the electrode surfaces of samples A and B during the first I-V measurements are available as supplementary material (video S1 and S2, respectively). The videos show the appearance of tiny flashes of light on the electrode surfaces during the measurement. FE-SEM images of sample A (Fig. 7(d) and 7(g)) and sample B (Fig. 7(e) and 7(h)) illustrate damaged areas generated on the electrode surfaces. These images show that a liquid-like material is ejected from a micron-sized hole, with numerous spherical particles around the damaged area: this phenomenon is similar to the explosion of dynamite buried under ground. The I-V curves of sample C indicate insulating characteristics in a stress voltage range of 0–20 V, and no damaged areas were observed on the electrode surface. The electric breakdown of sample A can be explained through the 1/E model.38,39 Electrical breakdown can result in two types of defects on the electrode surface, such as the worm-like and crater-like defects that are involved in the change of SiO2 breakdown characteristics upon Cu contamination.40 In this case, two possible processes explaining the operations of the devices are introduced in Fig. 8. For this hypothesis, a large amount of Sn submicron spheres embedded in the SiO2 layer may create a pathway for transferring electrons from the cathode to the anode. The large current flowing through narrowly connected parts of adjacent Sn spheres results in rapid Joule heating, which leads to the melting of the metallic Sn as in Fig. 8(b) and (e). The sudden volume increase of Sn from the solid to liquid phase creates a very large stress in the SiO2 layer, which ultimately breaks the SiO2 surface layer, followed by the ejection of Sn liquid in order to release the stress. After the surface destruction, adjacent damaged holes can be connected to the electrode by the remaining Sn on their walls as in Fig. 8(c) and Fig.7(g). For a thin oxide device, these residual Sn may shorten the circuit that results in the decrease of the breakdown voltage, as in the I-V curve of sample A (Fig. 7(a)). If we carefully observe the I-V curves of samples A and


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B (Fig. 7(a) and (b)), the initial current of sample A before the electrical breakdown is around 1010

A, whereas that of sample B before the electrical shutdown is from 10-3 to 10-2 A. It suggests

that there are many connected Sn spheres in sample B, while the Sn spheres are not joined to each other in sample A. After being heated by the large current, all connected Sn spheres are melt causing a bigger stress under the SiO2 surface of sample B comparing with that of sample A. That huge stress may not only break the thin SiO2 layer and the Ni/Au electrode on the surface but also eject almost Sn liquid out of the sample. The remained Sn in the damaged holes may be appeared in a form of discontinuous spherical Sn as in Fig. 8(f) and Fig. 7(h). Consequently, the electrical signal is shut down when the initial pathways are all broken, as in the I-V curve of sample B (Fig. 7(b)). The distribution of the damaged areas on the electrode surface of sample B after applying 20 V for 30 s is shown in Fig. 9(a) and (b). The projected area was calculated from the closed polygon surrounding the damaged area in the FE-SEM image (the detailed data are shown in Figure S4), yielding values in the range of 4–750 µm2 for this sample. Similar distributions of damaged areas were obtained at five different locations. The total projected area at the electrode center is 27.01%, which is approximately 2% higher than the average of the total projected areas on the electrode corners. The influence of stress time and stress voltage on the distribution of damaged areas of sample B is illustrated in Fig. 9(c,d) and Fig. 9(e,f), respectively. FE-SEM and polygonextracted images corresponding to these tests are shown in Fig. S5 and Fig. S6. Based on these data, the number of the damaged region on the electrode surface of sample B can be controlled by stress voltage and time. In addition, time-to-explosion tests (stress voltage of 20 V, stress time of 120 s) were carried out using small (video S3) and large electrodes (videos S4 and S5). The recorded videos indicate that the explosions are repeatable and uniform over a large area.


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CONCLUSIONS A simple method for the fabrication of β-Sn submicron-spheres embedded in SiO2 was introduced. The high degree of evaporation and diffusion of Sn into the SiO2 film results in the formation of β-Sn submicron spheres suspended in a SiO2 structure. By adjusting the working pressure, treatment time, and temperature, the thickness of the SiO2 film could be controlled from 300 nm to approximately 5700 nm. A possible mechanism for the material growth was introduced and discussed in detail. I-V measurements of sample A (fabricated at 700 oC and 10-1 Torr for 1 min) and B (fabricated at 700 oC and 10-1 Torr for 60 min) show different electrical characteristics, e.g., electrical breakdown and shutdown. During the measurements, tiny flashes of light were recorded on the electrode surface. The solid to liquid phase change of the Sn submicron sphere caused by Joule heating introduces a large amount of stress in the SiO2 layer, which ultimately leads to breakage of the SiO2 structure and of the electrode surface to release the stress. After the current-time measurements, the area of the damaged regions differs by just over 2% among five different sampling locations on the electrode surface. The surface damage can be controlled by adjusting the stress voltage (15–20 V) and stress time (1–120 s), and is independent of the size and shape of the electrodes. Due to the simplicity and integration capabilities of silicon fabrication technologies, the present materials may have great potential for application in self-destructing nanodevices.

METHODS Synthesis of the Sn nanostructures embedded in SiO2


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The Sn nano/microspheres embedded in SiO2 were grown on a single-crystal polished p-type Si (100) substrate (1 ohm·cm) by vacuum evaporation. The experimental setup is illustrated in Fig. S1(a) of the supplementary data. Commercial SnO powder (99.5% purity) was filled in an alumina boat and inserted into the center of a quartz tube (radius 20 mm, length 600 mm). The Si substrate (10 × 10 mm) was placed above the boat facing the SnO powder, with a distance of 8 mm between the Si substrate and the SnO powder. The growth was carried out in a vacuum horizontal furnace with operating temperature, pressure, and treatment time set in the ranges of 600–900 oC, 760–0.001 Torr, and 1–600 min, respectively. The temperature was increased at a rate of 5 oC/min, and the base pressure before the introduction of N2 was 0.0001 Torr. During the growth, ultrahigh purity nitrogen was flown into the system at a flow rate of 20 sccm. After treatment, the samples were cooled naturally to room temperature. The Si substrates were annealed in N2 (20 sccm) at 600–900 oC/0.1 Torr for 24 h without SnO powder to check the influence of residual oxygen (from the quartz tube or from the gas leakage) to the formation of the SiO2 layer. Cross-sectional FE-SEM images of the annealed Si substrates in Fig. S2 show that a thin layer of SiO2 (approximately 145 nm) was created on the surface of the Si substrate after annealing at 900 oC. For lower annealing temperature, the influence of the residual oxygen can be ignored.

SiO2 and β-Sn etching process β-Sn was removed from the SiO2 surface by sonication in HNO3 solution for 30 min. To etch both SiO2 film and β-Sn from the Si substrate, the sample was dipped in buffered oxide etch (BOE) solution for 40 min and HNO3 solution for 30 min. After treatment, all samples were cleaned by ethanol and dried by N2.


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Structure and morphology characterization of the SiO2 films embedded with β-Sn nano/microspheres The surface morphology of all samples and the cross-sectional structure of selected samples were characterized by FE-SEM (JSM-6701F) without conductive coating layer at the accelerating voltage of 10 kV. Structural, cross-sectional, and elemental mapping images of the sample synthesized at 700 oC and 0.1 Torr for 1 min were obtained by TEM (Telnai G2F20 STWIN, Philips). SPM data for the Sn/SiO2-stripped sample were obtained by a scanning probe microscope (Nanoscope IIIA, multi-mode system in tapping mode, Digital Instrument Inc.).

Electrical characterization of the SiO2 films embedded with β-Sn nano/microspheres Au/Ni (60/40 nm) electrodes were patterned on several selected samples by e-beam evaporation. A photograph and a scheme of the device are shown in Fig. S1(b). The I-V and current-time curves of the devices were recorded on a probe station (MSTech) with a semiconductor characterization system (Keithley 4200-SCS).


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Figure 1. Field emission scanning electron microscopy (FE-SEM) surface images of Si substrates treated at 700 oC for 1 h in N2 with working pressures of 760 Torr (a), 50 Torr (b), 5 Torr (c), 1 Torr (d), 0.1 Torr (e), and 0.001 Torr (f). Top-view and cross-section images of Si substrates treated at 700 oC for 1 min (g), 1 h (h), 5 h (i), and 10 h (j) in N2 with a working pressure of 0.1 Torr. Top-view and cross-section images of Si substrates treated at 600 oC (k), 700 oC (l), 800 oC (m), and 900 oC (n) for 1 h with a working pressure of 0.1 Torr. The scale bar is 1 µm.


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Figure 2. Elemental mapping images of the Si substrates treated at 700 oC and 1 Torr for 1 h in N2 (a–d). The scale bar is 1 µm.


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Figure 3. Transmission electron microscopy (TEM) image of the Si substrates treated at 700 oC and 1 Torr for 1 h in N2. The images in the insets highlight the β-Sn phase of the spheres embedded in the SiO2 layer (a). High-resolution transmission electron microscopy (HR-TEM) images of the interface between Si and SiO2 (b). Calculated fast Fourier transform (FFT) image (c) and interplanar spacing (d) and (e) indicating the presence of β-Sn in the Si layer below the interface. The scale bars in panels (a) and (b) are 100 and 5 nm, respectively.


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Figure 4. Field emission scanning electron microscopy (FE-SEM) images of samples treated at 800 oC and 0.1 Torr for 1 h (a, b). As-prepared sample after sonication in HNO3 solution for 30 min (c, d). As-prepared sample after stripping in buffered oxide etch (BOE) for 40 min and HNO3 solution for 30 min (e, f).


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Figure 5. Field emission scanning electron microscopy (FE-SEM) images of the Si substrates treated at 600 oC (a), 700 oC (b), and 800 oC (c) after etching in buffered oxide etch (BOE) for 20 min and subsequently in HNO3 for 40 min. 3D scanning probe microscopy (SPM) images of the corresponding samples treated at 600 oC (d), 700 oC (e), and 800 oC (f). Depth profile images of the samples along the red lines shown in the 3D images (g-i). The scale bar is 100 nm.


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Figure 6. Gradual formation of the β-Sn nano/microspheres embedded in the SiO2 layer (a–d).


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Figure 7. Current-voltage (I-V) curves for sample A (T = 700 oC, p = 10-1 Torr, t = 1 min) (a), sample B (T = 700 oC, p = 10-1 Torr, t = 60 min) (b), and sample C (T = 700 oC, p = 10-3 Torr, t = 60 min) (c). Field emission scanning electron microscopy (FE-SEM) images of the surface of samples A (d, g), B (e, h), and C (f, k) after the I-V measurements. The scale bars in panels (d-f) and (g-i) are 10 µm and 1 µm, respectively. The I-V curve measurements were repeated five times for each sample.


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Figure 8. Two possible processes for explaining the electrical breakdown in sample A (a–c) and the electrical shutdown in sample B (d–f) during the stress of bias voltage.


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Figure 9. Distributions of damaged areas extracted from the field emission scanning electron microscopy (FE-SEM) images of the sample fabricated at 700 oC and 0.1 Torr for 1 h, after timedependent tests with different electrode locations (a, b), stress times (c, d), and stress voltages (e, f).


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ASSOCIATED CONTENT Supporting Information available: − Schematic experimental setup of the chemical vapor deposition (CVD) system, 3D model of the device and Photographs of actual devices. − Cross-sectional FE-SEM images of Si substrates which were annealed in N2 at 600-900 o

C/0.1 Torr for 24 h without SnO powder.

− Energy-dispersive X-ray (EDX) spectrum of an area near a sphere-like grain and Elemental line scan data of Sn microspheres. − Distribution of damaged areas at different locations on the electrode surface. − Distribution of damaged areas on the electrode surface measured for stress times between 5 and 120 s. − Distribution of damaged areas on the electrode surface measured for stress voltages between 15 and 20 V. − Recorded videos of the electrode surfaces of samples A (treated at 700 oC and 10-1 Torr for 1 min) and B (treated at 700 oC and 10-1 Torr for 60 min) during the first I-V measurements. − Recorded videos of a small and large electrode surfaces of samples B (treated at 700 oC and 10-1 Torr for 60 min) during the time-to-explosion tests. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author * Young-Woo Heo. Tel.: +82 53 950 7587; Fax: +82 53 950 5465, E-mail: [email protected].


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Author Contributions 1

Vu Xuan Hien: designed and performed the experiment and wrote the paper.

2

Young-Woo Heo: designed and contributed to the experiment.

Funding Sources Authors Vu Xuan Hien and Young-Woo Heo received funding from National Research Foundation of Korea (NRF) Grants 2009-0093819 and NRF-2014R1A2A2A04005614.

ACKNOWLEDGMENT This work was supported by grants from the Priority Research Centers Program and the Midcareer Researcher Program through the National Research Foundation of Korea (NRF) funded by the Korean government (MSIP) (Grants. 2009-0093819 and NRF-2014R1A2A2A04005614).

ABBREVIATIONS PL, Photoluminescence; EDX, Energy-dispersive X-ray spectroscopy; ICCD, International Center for Diffraction Data; TEM, Transmission Electron Microscopy; HR-TEM, HighResolution Transmission Electron Microscopy; FFT, Fast Fourier Transform; FE-SEM, Field Emission Scanning Electron Microscopy; SPM, Scanning Probe Microscopy; I-V, CurrentVoltage; BOE, Buffered Oxide Etch.

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Table of Contents (TOC) graphic

A thick SiO2 film embedded with β-Sn-rich submicron spheres fabricated by a simple vacuum evaporation method is found to exhibit a microscale surface-destructive behavior upon application of a bias voltage. Substantial evaporation and diffusion of Sn into the SiO2 film results in the formation of β-Sn submicron spheres suspended in the SiO2 structure. During the I-V measurements, tiny flashes of light are recorded on the electrode surface. The I-V curves of the samples show both electrical breakdown and electrical shutdown, which result in damaged areas on the electrode surface. The damaged areas differ by only 2% among five different locations on the electrode surface. The explosion can be controlled by stress voltage (15–20 V) and stress time (1–120 s) and is independent on the size and shape of the electrode.


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Sn Spheres Embedded in a SiO2 Matrix: Synthesis and Potential

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