Crystallization of Si Templates of Controlled Shape, Size, and

Mar 23, 2015 - CNRS-Laboratoire de Photonique et de Nanostructures, route de Nozay, F-91460 Marcoussis, France. Cryst. Growth Des. , 2015, 15 (5), ...
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Crystallization of Si Templates of Controlled Shape, Size, and Orientation: Toward Micro- and Nanosubstrates Yann Cohin,*,†,‡ Frank Glas,‡ Andrea Cattoni,‡ Sophie Bouchoule,‡ Olivia Mauguin,‡ Ludovic Largeau,‡ Gilles Patriarche,‡ Elin Søndergård,† and Jean-Christophe Harmand*,‡ †

Surface du Verre et Interfaces, UMR 125 CNRS/Saint-Gobain, 39 quai Lucien Lefranc F-93303, Aubervilliers Cedex, France CNRS-Laboratoire de Photonique et de Nanostructures, route de Nozay, F-91460 Marcoussis, France

‡

S Supporting Information *

ABSTRACT: Fiber-textured polycrystalline silicon thin films have demonstrated their interest as seed layers for the epitaxial growth of high-quality materials on substrates such as glass or plastics. In the present work, we report a comprehensive study of the aluminum-induced crystallization (AIC) of isolated Si domains. Our study not only demonstrates the fabrication by AIC of shape- and size-controlled Si monocrystals at the nanoscale but also allows for the prediction of minimal annealing conditions for complete crystallization of such structures. These ultrathin [111]-oriented monocrystals are promising candidates as selected-area epitaxy substrates with both an easy control on the morphology of these structures and a low-temperature fabrication process. initially in contact with the a-Si film. Upon moderate temperature annealing, Si atoms, which are in contact with the metal, detach from the a-Si layer and diffuse into the grain boundaries of the Al film. This dissolution progresses until the concentration of mobile Si atoms in the polycrystalline Al reaches supersaturation,13 provoking the nucleation of Si crystalline grains. Under proper conditions, the Si nuclei adopt a specific crystallographic facet parallel to the interface of the Al film or to the substrate,14,15 which minimizes the interface energy.13,16 The growth of the Si grains continues by lateral homoepitaxy,17,18 until the a-Si is fully consumed, thereby causing the complete inversion of the Al and a-Si layers. This mechanism is called aluminum-induced layer exchange (ALILE).12 The polycrystalline film so obtained is composed of monocrystalline Si grains of several micrometers width. During the AIC of a-Si, the final grain size is set by the density of nuclei which form at the very beginning of the process. The formation of Si crystalline nuclei and their lateral extension consume the surrounding mobile Si atoms. As a consequence, the Si concentration in the Al layer decreases locally.19 Because of the high mobility of the Si atoms in the Al grain boundaries (in comparison, the dissolution of Si from the a-Si reservoir into the Al layer is a much slower process), the depletion region extends over a long distance (several micrometers) compared to the Al/a-Si stack thickness (tens to hundreds of nanometers).13,20,21 The presence of this depletion region around an existing nucleus suppresses new

U

sing a large-grained polycrystalline seed layer is a promising way to achieve the epitaxial growth of semiconductor materials on top of a low-cost foreign substrate, such as glass, metal, or plastics. The flexibility in the choice of foreign substrate may bring useful properties for many applications (transparency for sensors, flexible substrates1 for photovoltaics2−7 or electronics,8,9 etc.). To this aim, the control of the crystalline orientation of the seed layer is crucial: as an example, the growth of III−V semiconductor nanowires mostly requires a (111) surface.10 A step forward would be to implement this technology with patterned seed layers consisting of individual patches controlled in size, shape, and position. This would pave the way to material and device integration at precise locations on low-cost substrates. Moreover, defining seed layer domains of limited lateral sizes presents a great advantage: each domain can be monocrystalline. These perspectives lead us to propose the concept of monocrystalline micro- or nanosubstrates for epitaxy. In the following, we demonstrate how such a substrate can be achieved. Silicon is a good candidate to form seed layers on top of which various semiconductors can be grown. Generally, the crystallization of an amorphous Si (a-Si) layer is a high temperature process, which cannot be withstood by many other materials and which does not lead to oriented crystalline grains. Nonetheless, the aluminum-induced crystallization (AIC)11,12 of a-Si is an attractive method that can provide fiber-textured Si polycrystals at low temperature (from 150 to 550 °C). The substrate normal is the preferential direction toward which all the crystalline grains tend to align a common crystallographic axis. This method only requires a thin polycrystalline Al layer © XXXX American Chemical Society

Received: November 11, 2014 Revised: March 6, 2015

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Figure 1. (a) Schematics of the Al/a-Si bilayer pattern fabrication process, starting from thin films deposited on an oxidized Si(001) wafer, and using a dry-etching step. (b) Optical micrograph of a sample exhibiting the different square matrices. The gap between squares is 50 nm.

nucleations over a certain distance. The final shapes and sizes of adjacent Si grains are then directly inherited from the initial distribution of nuclei and from the extent of the depletion region. In the present work, we demonstrate that it is possible to fabricate patterns of isolated monocrystalline Si platelets with a [111]-orientation, thanks to the nucleation suppression mechanism described above. A comprehensive study of the limits of this method is pursued, as well as a complete evaluation of the nucleation kinetics of the isolated monocrystals. We finally show that dividing an Al/a-Si stack into small domains does not impact the AIC process and allows one to predetermine the localization of the Si monocrystals. Very thin films of Al (10 nm) and a-Si (15 nm) were deposited by magnetron sputtering on a thermally oxidized Si(001) wafer (25 nm-thick SiO2 layer). An intermediate oxide layer (SiOx) was intentionally deposited in the same chamber without breaking the vacuum (see Supporting Information). This oxide layer plays a critical role in the AIC mechanism. In particular, its presence favors a strong fiber texture of the final polycrystalline film. Both [111]15,17,27−31 and [001]14,15,28,29,31 fiber texture orientation can be obtained, depending on the oxidation state of this intermediate layer.14,22,29,31 We used conditions that are favorable to get a strong [111]-orientation of the grains after the AIC process. This Al/a-Si stack including the oxide interfacial layer was patterned by e-beam lithography and reactive ion etching (Figure 1a). Matrices of 20 × 20 squares, with side lengths ranging between 170 nm to 7.5 μm, were defined (Figure 1b). The squares consist of the Al/Si bilayer and are separated by 50 nm trenches where the bilayer was completely etched. The patterned substrates were annealed for various durations under nitrogen atmosphere at temperatures between 200 and 500 °C. Each annealing was performed at constant temperature. The resulting AIC Si layers present the best surface coverage for the lowest annealing temperature (200 °C). At higher annealing temperature, the fiber texture quality is improved, but the Si layer morphology presents dendrites and surface coverage is less uniform. To optimize both the surface coverage and the fiber texture, two-step annealing consisting of 10 h at 200 °C and 10 min at 500 °C is efficient. However, since the present study focuses on kinetic aspects of AIC, we did not use such a two-step method which would have led to a more complex interpretation of our experimental results. Then, the samples were etched in a 5% HF solution to remove the Al top layer as well as the AlSiOx interlayer23 before their observation by scanning electron microscopy (SEM) and tapping-mode atomic force microscopy (TM-AFM). In AIC, when the conditions are appropriate, the Si grains grow from crystalline nuclei by lateral homoepitaxy, until either they get in contact with each other or until all the available a-Si is consumed. Each of these grains is monocrystalline since the orientation of the nucleus is maintained during the lateral

growth.22 In a first series of experiments, we studied the influence of the temperature on the AIC of the Al/a-Si domains at long annealing times. The resulting crystallization state was observed on the series of patterned squares as well as on nonpatterned areas of the same sample (Figure 2a). On the

Figure 2. (a) SEM image of an AIC Si polycrystalline layer in a nonpatterned area: the hillocks are visible, as well as holes between the different monocrystalline grains. (b) Morphology of holes and hillocks observed by AFM. (c) Profile across holes and hillocks: three distinct height levels with 10 nm steps are observed.

latter, we notice an incomplete percolation of the Si grains, with holes at their frontiers. This is related to the particular Al/a-Si ratio chosen for this experiment (10 nm/15 nm), which is not perfectly optimized for forming a continuous polycrystalline Si layer. A slightly thicker Si layer would be necessary to complete the layer (see Supporting Information, Figure S4). The holes originate from wet-etched Al grains that did not participate in the crystallization and which are therefore sitting directly on the SiOx substrate. These holes draw a frontier surrounding each individual Si monocrystal (Figure 2a). They are very useful to distinguish and count the monocrystalline Si grains. As shown B

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Crystal Growth & Design by TM-AFM, the crystallized Si grains adopt two distinct heights of 10 and 20 nm (Figure 2b,c).22,24,25,32 Other authors showed that this roughness originates from a second level of crystallization of remaining a-Si, induced by Al grains that already exchanged their position with the first level of crystallized Si. Hence, the fact that the crystalline Si seems to increase in height by steps of 10 nm must be related to the size of the Al grains which are replaced by Si during the layer inversion. In our particular case, the Al layer is also 10 nm high, and most probably, the Al grains have the same height. Therefore, a height of 20 nm should correspond locally to two successive Al/Si exchanges involving the same Al grain. The roughness produced by such events can be reduced by adjusting the Al/a-Si ratio and is not detrimental to our demonstration. However, this roughness might become an obstacle in the prospect of using the platelets for epitaxial growth. A specific surface treatment might be required to reduce the roughness if the sole adjustment of the initial thicknesses proves to be insufficient. To summarize the effect of the Al/a-Si thickness ratio, we observed that the crystallized Si layer is incomplete at a thickness ratio lower than 1.5 and that a second level of crystallized Si starts for a thickness ratio higher than 1. Hence, the conditions used in our experiments correspond to a situation where the surface morphology of the crystallized AIC samples presents three possible height levels. These height levels produce different intensities in the SEM images (Figures 2 and 3): level 0 (bright) corresponds to the SiOx regions that are not covered by crystalline Si, levels 1 (gray) and 2 (dark) correspond to the regions where the Si grains are about 10 and 20 nm high, respectively. The SEM images were processed in order to delineate the Si grains. The method relies on the definition of a threshold intensity set between intensities of levels 0 and 1. This threshold is used to obtain binary maps from the SEM images. Binary maps were obtained by applying a contrast threshold to the images. By eroding the black areas representing the monocrystalline grains, we could easily isolate and count them. This counting procedure was implemented on the patterned samples (Figure 3a): the average number of monocrystals per crystallized square domain is shown in Figure 3b as a function of the square side length L and of the annealing temperature T. For the lowest temperature (250 °C), a fraction of the squares did not crystallize (this is clearly seen in the squares of sides smaller than 220 nm). This phenomenon, explained later in this work, is taken into account in the graph of Figure 3b, where we restrict ourselves to the analysis of crystallized domains. The graph exhibits two regimes. For small L, a single Si monocrystalline grain per square was formed (Figure 3c−e), while for larger L, several grains per square are observed (Figure 3f,g). Note that in the SEM images shown in Figure 3c−e, the contrast visible on the single grains is due to regions at level 1 or level 2. In Figure 3f,g, grain boundaries (level 0) appear with a brighter intensity. The number of grains per square increases with L to tend asymptotically to values corresponding to the grain surface density dS measured on the nonpatterned area (the asymptotes for different temperatures are plotted as straight dashed lines in Figure 3b). The transition from mono- to polycrystals per square depends on the crystallization temperature: a higher temperature leads to smaller grains (higher surface density) and thus to a transition at smaller square side lengths. With the present conditions, the largest monocrystalline squares that we can fabricate are about

Figure 3. (a) Illustration of the grain counting method: (1) a binary threshold is applied to a SEM image (here an example of four adjacent squares), (2) the binary image is eroded in order to isolate the monocrystals; (3) the grains are counted. (b) Number of monocrystals per square as a function of the square side length. Insets (c−g) show SEM images of representative squares with initial side lengths of 170 nm, 220 nm, 770 nm, 3.0 μm, and 5.0 μm.

450 nm wide at 500 °C, 700 nm wide at 350 °C and 1 μm wide at 300 and 250 °C. The surface of the squares was further characterized by TMAFM. Figure 4a shows a section of the 170 nm wide square array: most of the monocrystalline Si domains have a relatively smooth morphology (Figure 4b), although a few of them exhibit more irregularities at the surface or near the edges (the typical height standard deviation is 0.7 nm). The morphology of the Si squares reproduces likely the initial Al grain size distribution and its roughness.26 Indeed, the lateral size of the hillocks is comparable to the typical Al grain size (20−30 nm). This is clearly visible in the phase signal of the TM-AFM image of Figure 4c, and it confirms that Si has filled the volume that was occupied by individual Al grains before the AIC. The crystalline orientation of large patterned arrays was characterized by grazing incidence X-ray diffraction (GIXRD) (Figure 5). To this aim, a 1 × 1 cm2 array patterned with 730 nm side squares was annealed in two steps: a first step at 300 °C for 10 h to ensure that each square consists of a single grain, and a second step at 500 °C for 10 min to improve the crystalline quality of the Si patches. We recall that the initial stack is optimized to produce a [111] fiber texture. On this sample, the Al layer was not removed. The diffractogram exhibits two series of peaks originating from the Al grains and from the Si crystals respectively (Figure 5). The peaks related C

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which are due to twin planes parallel to the substrate,22,33−35 and traces of Si {111} and Si {311} orientations. The dominance of Si {220} and Si {422} is the signature of a strong [111] fiber texture of the Si grains, while the minor peaks can be attributed to the defective orientation of a small fraction of the grains. The volume of this small fraction can be estimated by comparing the areas under the different Si diffraction peaks divided by their own structure factor. This estimation leads to a fraction of misoriented grains of 7%. A similar experiment performed on a nonpatterned sample fabricated in the same conditions gave comparable results. Rotating the sample along the ϕ axis does not modify the intensity of the diffraction peaks. This means that the in-plane orientation of the Si grains is homogeneously distributed around this axis and hence that the square sides do not influence the in-plane orientation of the Si crystals. Hence, the addition of a lithography step to the AIC enables the fabrication of shape- and size-controlled polycrystalline Si platelets. These patterned platelets present a low roughness and an excellent alignment of their [111] crystalline axis along the normal direction. Moreover, we determined the critical lateral size of the patterns below which each patch consists of a monocrystal. These monocrystalline patches could serve as micro- or nanosubstrates for epitaxial growth on various supports. We shall now explore the kinetics of the AIC in patterned Al/a-Si samples. Nucleation and subsequent lateral growth can be investigated in more detail by varying the annealing time. In particular, nucleation and lateral growth rates are accessible in patterned samples by observing the crystallization of an array of numerous Si patches at shorter annealing times. Since we observed strong similarities between the AIC of a-Si on nonpatterned and patterned areas, we start by studying the nucleation and crystallization kinetics of a full-wafer Al/a-Si stack to extract relevant quantities in this reference system. The samples were annealed at different temperatures for various durations. The grains were characterized in terms of grain surface density dS and mean grain area A, by processing SEM images. The Johnson−Mehl−Avrami−Kolmogorov (JMAK) phase transformation theory36 was previously used by other authors37 to model the nucleation and growth in two dimensions of Ge crystals obtained by AIC. In the present work, we model the evolution of grain surface density (Figure 6a) as a function of time by considering a nucleation rate per unit surface and unit time α independent of time and location.

Figure 4. (a) Tapping mode atomic force micrographs (TM-AFM) of an array of 170 nm-wide Si squares. (b, c) TM-AFM images ((b) height and (c) phase) of four adjacent squares selected in (a). The scans were made at 45° in order to minimize artifacts; the images were reoriented in post-treatment. The roughness of the Si platelets, inherited from the Al initial grains, is visible.

Figure 5. Grazing incidence X-ray diffractogram of an annealed 1 × 1 cm2 sample patterned with 730 nm side squares. The inset illustrates the GIXRD configuration employed here.

to the Al layer (Al {111}, {200}, and {220}) show a powderlike polycrystal. The second series is related to Si: it consists of intense Si {220} and {422} peaks, extremely weak Si 1/3{422}

Figure 6. AIC crystallization kinetics of self-organized Si grains on a full wafer. (a) Surface density dS as a function of annealing time and temperature. (b) Mean equivalent radius R of the grains as a function of annealing time and temperature. (c) Schematics of an ensemble of N growing grains (black discs) having a mean radius R(t) and depleting a larger surface, whose equivalent mean radius is R(t) + Rdep. D

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Crystal Growth & Design Table 1. Surface Density dS, Radial Growth Velocity vr, Nucleation Rate α, and Depletion Radius Rdep Obtained from Experiments on Full-Wafers at Different Temperatures T (°C) 200 225 250 275 300

surface density, dS (μm−2) 0.17 0.20 0.26 0.31 0.37

(±0.02) (±0.02) (±0.02) (±0.02) (±0.02)

radial growth velocity, vr (μm s−1) −4

1.82 × 10 (±0.06 5.62 × 10−4 (±0.08 1.8 × 10−3 (±0.2 × 4.5 × 10−3 (±0.1 × 1.2 × 10−2 (±0.1 ×

−4

× 10 ) × 10−4) 10−3) 10−3) 10−2)

We define the mean equivalent radius R of the grains by A = πR2. Both dS and R are plotted as functions of the annealing time in Figure 6a,b at each temperature. The equivalent radius R varies linearly with annealing duration (Figure 6b), before it saturates as the grains complete the film. During the linear radial growth regime, R verifies: R = νr(t − t0), where vr (μm s−1) is the radial growth velocity and t0 an incubation time related to sample thermalization and to the dissolution time of a-Si before supersaturation in the Al layer.18,30,38,39 The nucleation probability to nucleate dN grains per time dt is αS dt, where S is the available surface for nucleation. Assuming that N grains have nucleated on the total surface S0 at an annealing time t (schematized in Figure 6c) and that these grains have a mean radius R(t) and a constant depletion radius Rdep, the available surface S(t) is S0 − π(Rdep + R(t))2N(t). Considering that N(t) = S0dS(t), the differential equation that rules the grain nucleations is

nucleation rate, α (μm−2 s−1) 1.1 5.4 2.7 8.5 2.2

× × × × ×

10−4 10−4 10−3 10−3 10−2

(±0.1 (±0.1 (±0.2 (±0.1 (±0.2

× × × × ×

10−3) 10−3) 10−3) 10−3) 10−3)

depletion radius, Rdep (μm) 0.82 0.83 0.75 0.66 0.54

(±0.05) (±0.05) (±0.05) (±0.05) (±0.05)

the depletion radius is found to monotonously decrease with the temperature from 0.83 μm at 225 °C to 0.54 μm at 300 °C. Such a decrease was already observed by other authors.19,43 This could be related to the thermal activation of the dissolution of Si from the a-Si layer into the Al layer. A faster Si atoms supply to the Al layer could partially compensate the depletion by nucleation and crystal growth. The depletion radius values are fully consistent with the transition between the two regimes of Figure 3b for the patterned squares. Building on the comprehensive kinetic and thermodynamic study carried out on full wafer samples, we now characterize the crystallization of the Al/a-Si squares. We annealed a patterned substrate for 5.4 × 103 s (1.5 h) at 250 °C, i.e., in conditions that lead to a complete crystallization of the full wafer samples (Figure 6). We consider only the squares that are sufficiently small to contain a single monocrystal. After HF-etching, the crystallized squares can easily be distinguished from those that remained amorphous (Figure 8a) because of their strong SEM contrast (Figure 8b). It appears clearly that this annealing procedure is not sufficient to crystallize every square, as evidenced on Figure 8c, which shows crystallized squares and squares that are still amorphous. However, the square side length has a very strong effect on the probability of nucleation (all other things being equal) (Figure 8d): smaller squares have a lower probability to nucleate than bigger ones. The fabrication with certainty of crystallized Si squares thus demands a more precise study of the nucleation in patterned samples, considering all relevant parameters, namely, annealing time t and temperature T, but also square size L. To this end, we plot the nucleated fraction φ, i.e., the ratio of the number of nucleated squares N to the total number N0 of squares, as functions of the square size L (Figure 9a) and of the annealing time t (Figure 9b). In order to explain the variations of φ with t and L, we first consider, as we did in the nonpatterned case, a nucleation rate per unit surface and unit time β, independent of time and location. The probability dP to form a grain per unit time dt is equal to the product of β by the surface of a square L2: dP = βL2 dt. Nucleation can only happen in the non-nucleated (N0− N) squares, since only one monocrystal can nucleate on each single square (as seen in Figure 3), because the area of the squares considered is smaller than the depleted area; thus dN = β(N0−N)L2 dt, which can be rewritten as

dd S (t ) = α(1 − π (R dep + vr(t − t0))2 dS(t )) dt

First, vr and t0 are determined from the evolution of R as a function of annealing time at each temperature (Figure 6b). Considering these values, α and Rdep are obtained by minimizing the error between the experimental data and dS(t), which was calculated by solving the previous differential equation via a first-order Runge−Kutta method (solid lines in Figure 6a). The nucleation rate α obtained (Table 1) is mainly determined by the increasing part of dS(t), before the depletion regions overlap. The evolution with annealing temperature of the nucleation rate α (Figure 7a) and radial growth velocity vr (Figure 7b) are then analyzed. As expected, these two parameters increase with temperature.21,27,40 Arrhenius plots of the extracted values give activation energies. The activation energies for nucleation and growth, respectively 1.1 (±0.1) eV and 0.9 (±0.1) eV, are close to each other in agreement with other studies.21,27,41,42 Besides,

∂φ (t ) = (1 − φ)βL2 ∂t L

The solution of this first order differential equation with φ(t 2 = 0,L) = 0 is φ(t,L) = 1 − e−βL t, to which a delay t0 may be applied as suggested by the experimental data (Figure 9b). Even if this simple model adapts well to the shape of the data of Figure 9, it cannot explain some of their features: on the one hand, we observe a shift in L (as if the surfaces of the squares

Figure 7. (a) Nucleation rate α as a function of 1000/T on a semilogarithmic scale (Arrhenius plot). (b) Arrhenius plot of the radial growth rate vr. E

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Figure 8. (a) After a given time, every Al−a-Si square may not have crystallized. The crystallized squares can be revealed by removing their top Al layer by wet etching. (b) SEM images of HF-etched crystalline-silicon (left) and amorphous-silicon (right) squares. (c) SEM image of a 20 × 20 matrix of 550 nm wide crystallized (bright) and still amorphous (dark) squares, obtained after etching a sample annealed at 250 °C for 6 × 103 s. (d) Maps of the crystallization state of six arrays with different square sizes, annealed for 5.4 × 103 s at 250 °C.

Figure 9. Nucleated fraction φ as a function of the square side length L for different annealing times (a) at 250 °C and as a function of annealing time at different lengths (b): experiments (symbols) and nucleation model (dashed lines). (c) Schematics of a damaged square: the red area cannot nucleate. (d) and (e) Gaussian distributed damaging, fitted from the data at 250 °C: 91% of the 220 nm wide squares are expected to nucleate (d) and only 65% of the 170 nm wide ones (e).

were overestimated); on the other hand, for the smaller pattern sizes (170, 220, and 350 nm) φ does not tend to 1 at long annealing times, at variance with our model. Note that these imperfect fits are not shown in Figure 9. The overestimation of the effective square domain area can be attributed to some damaging of the square edges during the sample preparation: the amorphization of the edges of the Al films over a width L0 during ion etching may prevent its participation in the AIC of the a-Si squares (Figure 9c). However, replacing L2 by (L − L0)2 in the differential equation does not modify the limit of φ at long annealing time. Indeed, either the damaged area is larger than the square side length, and then, no nucleation is expected: φ(t,L < L0) = 0; or else L > L0, and the nucleation should occur after a sufficiently long

and the Gaussian damaging distribution, defined by its center L0, and its standard deviation σ (the model equation is given in the Supporting Information as well as a simpler analytical expression of φ (t, L)). The three parameters (β, L0, σ) were obtained through a fit of the experimental data at 250 °C. The agreement is excellent (dashed lines in Figure 9a,b). Similar nucleation experiments were carried out at 300 and 350 °C (data and fits are shown in the Supporting Information). Table 2 summarizes the results obtained from the fits at the three temperatures. At 250 °C, as expected from the large fraction of squares that are unable to nucleate for the smaller domain sizes (Figure 9b), the average damaged width is large (146 nm) and broadly distributed (56 nm) (Table 2). The plot of this Gaussian

2

annealing time, according to φ(t,L > L0) = 1 − e−β(L−L0) (t−t0). This issue was addressed by considering that the width of the damaged area is not the same for each square of a given side length, but has a Gaussian distribution, which means that the equivalent active surface of the squares varies from one domain to another. In this model, each square has its own effective surface and, depending on the parameters of the Gaussian distribution, some squares may be totally damaged and never crystallize, whatever the annealing time. We implemented a modified nucleation model combining the model previously exposed, for which the sole parameter is the nucleation rate β,

Table 2. Nucleation Rate β and Gaussian-Distributed Damaging Parameters (L0, σ) Obtained from Experiments at Different Temperatures nucleation rate

F

Gaussian-distributed damaging

T (°C)

β (μm−2 s−1)

L0 (nm)

σ (nm)

250 300 350

2.6 × 10−3 (±0.05 × 10−3) 1.6 × 10−2 (±0.05 × 10−2) 2.1 × 10−1 (±0.05 × 10−1)

146 (±6) 56 (±6) 61 (±6)

56 (±5) 4 (±2) 4 (±2)

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lithography step to their technologies. On the other hand, by simply increasing the temperature from 250 to 300 °C, it is possible to reduce drastically the damaging of the domain edges, which gives access to smaller crystallized patterns. An ion-etching-free lithography, such as a lift-off process, can also be implemented to reduce the defects at the pattern edges. The damaged region can be limited to a 30 nm crown which seems to correspond to the Al grain size and appears ineffective for crystallization at any temperature. The Si monocrystalline platelets can be used for various purposes, in particular, as seeds for selected area epitaxy. They can be precisely localized and their shape can be controlled. Their [111] orientation is well suited to semiconductor epitaxial growth, and their size can be tuned to obtain epitaxial nano- or microstructures, such as nanowires23,28,44 or larger pads of semiconductor with the idea that a single device could be fabricated on one pad. In this latter case, pads of a few μm2 would be suitable. These perspectives open for a novel concept of epitaxial micro- or nanosubstrates for semiconductor integration on foreign supports in a broad range of applications.

function (Figure 9d,e) shows that a large fraction of the smallest domains (170 and 220 nm) cannot nucleate. Indeed, only 65% of the 170 nm wide squares, and 91% of the 220 nm wide ones are not fully damaged. At higher temperatures (300 and 350 °C), the damage is drastically reduced: approximately 30 nm are removed from the edges (leading to a total damaging about 60 nm wide). We attribute this drop in the damaging to a recrystallization of the Al layer during the annealing procedure, which cures the amorphization by ion etching observed at 250 °C. Nevertheless, a slight reduction of the available active surface enabling nucleation is still observed. The 30 nm damaged crown actually corresponds in width to a single Al grain. Two main reasons could explain this similarity: the Al grains on the periphery are exposed to air (not being protected from the side by the a-Si layer) and can be fully oxidized; they can also be improper for nucleation because of geometrical reasons, such as the configuration of their grain boundaries at the domain edges. In such a case, the width of the ineffective crown could be shrunk by reducing the Al grain size. The experiments on both types of samples gave access to a nucleation rate (α or β). We plotted α and β as functions of inverse temperature (Figure 10). The two series of data

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EXPERIMENTAL SECTION

Bilayer Deposition. Thin films of Al (10 nm) and a-Si (15 nm) were deposited by magnetron sputtering on a thermally oxidized Si (001) wafer (25 nm-thick SiO2 layer). An intermediate oxide layer was intentionally deposited in the same chamber without breaking the vacuum. This oxide layer was formed by preparing the Si target before the deposition. This target was exposed to an Ar/O2 gas mixture (30 and 6 sccm, respectively) in order to obtain a controlled oxide layer at its surface. For deposition, the magnetron sputtering apparatus was used at 100 W direct-current (DC) power for both Al and Si (with respective deposition rates of 0.435 nm s−1 and 0.212 nm s−1) under a pure Ar atmosphere of 1.5 μbar. Pattern Fabrication. (Figure 1a): On top of the Al/a-Si stack, a polymethyl methacrylate (PMMA) A5 film was spin-coated at 5000 rpm with an acceleration of 2000 rpm s−1 for 30 s. It was then cured for 15 min at 125 °C on a hot plate. Lines were written in the resist by electron-beam (e-beam) lithography using a Vistec EBPG 5+ apparatus, operated at 100 kV with a current of 5 nA and a dose of 1200 μC cm−2. On the same sample, we drew 12 matrices of 20 × 20 squares, with side lengths of 170, 220, 350, 420, 550, and 770 nm, as well as 1.3, 2.0, 3.0, 4.1, 5.0, and 7.5 μm, as illustrated in Figure 1b. The resist was developed in a methylisobutylketone/isopropanol (3:1) mixture for 40 s, rinsed in isopropanol, and dried with nitrogen. The squares of different sizes were transferred into the Al/a-Si stack down to the thermal SiO2 layer using inductively coupled plasma (ICP) reactive ion etching (RIE). A BCl3/HBr chemistry was used to ensure anisotropic etching of both Al and a-Si, with high selectivity against PMMA. A low ICP power and a low DC bias were also used to enhance the selectivity and to minimize lateral amorphization of the etched materials. The typical etching time was about 2 min. The lithography resist was then removed by dipping the sample into trichloroethylene at 90 °C for 5 min. The sample was then rinsed sequentially in thrichloroethylene, acetone, and isopropanol before being dried with nitrogen. Annealing and Al Removal. The patterned substrates were annealed under a nitrogen atmosphere at temperatures between 200 and 500 °C for various durations (from 1 min to 22 h). The samples were introduced in the oven after its temperature had stabilized. After annealing, the samples were etched in a 5% aqueous solution of HF to remove the Al top layer as well as the AlSiOx interlayer,24 which makes easier the scanning electron microscope (SEM) observations. Sample Series. Two series of annealing experiments were carried out on these samples: − Counting the number of monocrystals that compose each square of the patterns after annealing for 20 h at 250, 300, 350, or 500 °C.

Figure 10. Arrhenius plots of nucleation rates α and β. Gray symbols stand for the data obtained during the full-wafer experiments (α), while red symbols correspond to patterned samples (β).

perfectly align in the Arrhenius plot, which proves that the corresponding processes have the same prefactors, and the same activation energies. We therefore conclude that the nucleation of Si grains in the cases of the full-wafer Al/a-Si stack and of separated domains follows exactly the same process. This property was expected, since β is independent of the domain area. In this work, we demonstrated the compatibility between the AIC of amorphous Si and a patterning of the material by lithography and ion etching at various scales. Associating these two technologies enables the fabrication of [111]-oriented Si monocrystalline platelets. The crystallization of such monocrystals meets two limits: first, a monocrystalline platelet cannot exceed the lateral size of a Si grain grown on a nonpatterned sample (∼1 μm in our conditions); second, our experiments and modeling demonstrate a damaging of the periphery of the etched patterns, which limits the minimum Si platelet size. We firmly believe that these two limitations can be overcome. Indeed, on the one hand, by modifying the deposition methods, conditions, and procedures, other authors already obtained wider AIC Si grains (10 μm large and above) on nonpatterned wafers (however, they used thicker Al/a-Si stacks),14,31 and it seems straightforward to integrate the G

DOI: 10.1021/cg5016548 Cryst. Growth Des. XXXX, XXX, XXX−XXX

Article

Crystal Growth & Design − Studying the crystallization kinetics of a-Si by AIC on full-wafer and patterned substrates using anneals of various durations from 120 s to about 8 × 104 s (22 h). Sample Characterizations. SEM images were taken using an FEI Magellan 400 XHR apparatus operated at 10 kV and 0.2 nA. TM-AFM was performed using a Bruker Icon III equipped with Budget Sensors Al-coated tips. GIXRD was carried out using a Rigaku Smartlab with a Cu anode (λ = 1.540598 Å).

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ASSOCIATED CONTENT

S Supporting Information *

Complete model for nucleation kinetics. Additional results at various temperatures. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Authors

*(Y.C.) E-mail: [email protected]. *(J.-C.H.) E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS L.P.N. is a member of CNRS-RENATECH national network of large nanofabrication facilities.

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REFERENCES

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DOI: 10.1021/cg5016548 Cryst. Growth Des. XXXX, XXX, XXX−XXX