Silicon Nanomembranes with Hybrid Crystal Orientations and Strain

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Silicon Nanomembranes with Hybrid Crystal Orientations and Strain States Shelley Scott, Christoph Deneke, Deborah M Paskiewicz, Hyuk Ju Ryu, Angelo Malachias, Stefan Baunack, Oliver G. Schmidt, Donald E. Savage, Mark A. Eriksson, and Max G. Lagally ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b14291 • Publication Date (Web): 11 Nov 2017 Downloaded from http://pubs.acs.org on November 11, 2017

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ACS Applied Materials & Interfaces

Silicon Nanomembranes with Hybrid Crystal Orientations and Strain States Shelley A. Scott*, Christoph Deneke2,3,4, Deborah M. Paskiewicz1,*, Hyuk Ju Ryu1,#, Angelo Malachias5, Stefan Baunack3, Oliver G. Schmidt3, Donald E. Savage1, Mark A. Eriksson1, and Max G. Lagally1

AUTHOR ADDRESS 1 University of Wisconsin, Madison, Wisconsin 53706, USA 2 Laboratoria Nacional de Nanotechnologia, Centro Nacional de Pesquisa em Energia e Materiais, 13083-100, Campinas, Brazil 3 IFW Dresden, Helmholtzstr. 20, Dresden, Germany 4 Instituto de Física "Gleb Wataghin", Universidade Estadual de Campinas (Unicamp), 13083-859, Campinas, SP, Brazil 5 Universidade Federal de Minas Gerais, CP 702, 30123-970 Belo Horizonte

KEYWORDS Epitaxy, selective growth, hybrid crystalline materials, Silicon nanomembranes, Interfaces, Strain Engineering

ABSTRACT

Methods to integrate different crystal orientations, strain states, and compositions of semiconductors in a planar, and preferably flexible, configuration may enable non-traditional sensing, stimulating, or communication device applications. We combine crystalline-silicon nanomembranes, patterning, membrane transfer, and epitaxial growth to demonstrate planar

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arrays of different orientations and strain states of Si in a single membrane, which is then readily transferable to other substrates, including flexible supports. As examples, regions of Si(001) and Si(110) or strained Si(110) are combined to form a multi-component single substrate with highquality narrow interfaces. We perform extensive structural characterization of all interfaces and measure charge carrier mobilities in different regions of a 2D quilt. The method is readily extendable to include varying compositions or different classes of materials.

1. Introduction The integration of single-crystal semiconductor materials with different crystalline orientations, strain states, and/or compositions into a planar sheet configuration suggests a range of concepts for increasing device functionalities on a single, preferably flexible, substrate. Strain modifies the band structure,1,2 thereby also modifying carrier transport and mobility; crystallographic orientation influences carrier transport,3,4 and composition can influence optoelectronic properties and transport.5,2 Combining materials via epitaxial growth alone restricts the selection to those with same or very similar lattice constants. Bonding of regions of disparate materials carries the risk of formation of interface defects that may degrade performance of the composite system, and also makes planar structures difficult to envision and to realize. When, on the other hand, the integration involves the transfer and bonding of very thin sheets of single-crystal semiconductors, called nanomembranes (NMs), novel material combinations become possible6,7 and defect formation at the bond interface can be avoided, in


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contrast to the bonding approaches mentioned above.8,7 In general, NM transfer produces “vertical” integration, in that a sheet is added on top of a substrate, thus changing its function. If the NM is patterned, then laterally varying properties can be achieved, but at the expense of having a non-planar structure. In many cases a flat surface with laterally varying function would be desirable, especially for deformable, flexible devices. We describe here a facile method for creating such laterally varying function. The method involves the transfer of a patterned NM to a host that is not just a passive support, but whose functional properties are desired in the final structure, and combines the transfer with epitaxial growth to planarize the structure. The pattern in the NM can be any shape or dimension, and the NM can be any material that can be created in NM form. The host substrate can be rigid or itself fabricable as a NM. We use single-crystal silicon nanomembranes (SiNMs), generated from silicon-oninsulator (SOI) to demonstrate the method. SiNMs can be patterned and etched prior to transfer, using standard lithography techniques, and used as growth substrates. They form a high-quality bond when transferred and bonded to the same or dissimilar materials,7 and in terms of growth, have allowed heteroepitaxy of large-lattice-mismatch materials (such as Ge9 and InAs10,11 on Si) that cannot be achieved with growth on rigid substrates. SiNMs have already greatly expanded the use of Si to new applications.12,13,6 For hybrid-orientation Si sheets, we transfer Si(110) membranes, patterned and etched to create an array of holes, to Si(001) and SOI(001) hosts, and then epitaxially grow Si using chemical vapor deposition (CVD). Because CVD uses vapor phase transport of a precursor gas, the entire structure is accessible to growth, allowing epitaxy on the sidewalls of the holes


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patterned in the Si(110)NM as well as the SOI(001) host that is exposed by the holes in the NM. In this particular case, a known dependence of the CVD growth rate on Si substrate orientation [10X faster on Si(001) than on Si(110)] allows the entire structure to attain a planar geometry, i.e., the holes are “filled” without significant growth on the transferred (110) membrane. The result is a planar quilt with different surface orientations that is laterally essentially continuous, with narrow domain boundaries. Charge transport properties in the regions are representative of values in the corresponding bulk crystals. The method is extended to incorporate strained Si(110) by using elastic strain sharing amongst the layers of a temporarily freestanding Si/SiGe/Si trilayer NM that avoids altogether the generation of misfit dislocations,14,15 in place of the original Si(110)NM. In this case, the final structure is a planar quilt of regions of Si(001) and strained Si(110), a hybrid-strain NM, or hybrid strain/orientation NM. Additionally, compositionally laterally varying planar membrane structures [a hybridcomposition NM] can be created by transferring a NM of the desired composition (e.g., SiGe, Ge, GaAs) with a particular orientation or strain state to SOI and using epitaxy and liftoff, as further described later.16 In other words, a differential growth rate on the two regions of the quilt is not a requirement In all cases, the use of SOI as the host substrate [or, more generally, any multilayer structure containing an etchable release layer] allows a second transfer process, whereby the hybrid structure can be transferred to any other host, or even rendered free-standing, in a contiguous-membrane form.


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There are many potential applications of laterally heterogeneous planar crystalline sheets. By way of example, we measure the charge carrier mobility on different regions of a Si(110)/Si(001) hybrid-orientation quilt. In the traditional (001) orientation of Si, the hole mobility significantly lags the electron mobility. Holes have a lower effective mass in the Si(110) plane, enabling hole mobility gains of about twice that of the Si(001) orientation, but at the expense of a degraded electron mobility.17 Introducing tensile strain to Si(110) is known to give an additional boost to the hole mobility.18 Tensilely strained Si(110) grown on conventional strain-graded substrates suffers from a high density of misfit dislocations (an order of magnitude higher than Si(001) with an equivalent strain),19 which can negate the potential hole-mobility gains. Elastic strain sharing among the layers of multi-epitaxial-layer NMs described here avoids this difficulty and provides a platform for mitigating misfit dislocation formation and propagation.15,20 Consequently, mobilities can be optimized in strained-NM structures.

2. Results and Discussion 2.1 Fabrication of transferable hybrid-orientation nanomembrane systems Silicon NMs in their most general form are thin sheets of single-crystalline Si, with thickness ranging from 2 nm to >200 nm.14,21 Generally, they begin as the template layer of silicon on insulator (SOI), which is composed of a thin layer of Si (template) separated from a bulk Si(001) wafer (handle) by a buried oxide layer (BOX). They can be released from this original host substrate, generating a temporarily free-standing sheet, which can readily be bonded to a new host, of almost any other material. Similarly, semiconductor-on-insulator with a template layer of Si(110), Si(111), strained Si(001) or Si(110), Ge, or SiGe can be obtained or


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fabricated, and NMs of these and other materials can be grown using various epitaxial methods.6,16 The process sequence for fabricating (110)/(001) hybrid-orientation structures from NMs is schematically depicted in Figure 1. The procedure is similar for all the above-mentioned template layers and is thus inclusive of varying composition (which we will show later for a Si/SiGe/Si heterostructure) and varying strains. The BOX is selectively removed with wet chemical etching (see Methods), either directly or, for large areas, via the introduction of lithographically patterned etch holes, Figures 1b and 1c. Removal of the release layer (here the buried oxide) forms, for SOI(110), a temporarily free-standing Si(110)NM, which is either bonded back in place on the original Si(001) host (shown in Figure 1d), or transferred to a new SOI(001) substrate (or an entirely different host). The etchant access holes in the Si(110) NM also define the regions where the Si(001) crystal plane of the host is exposed. Si is deposited over the structure illustrated in Figure 1d using chemical vapor deposition (CVD). Growth of Si via CVD on Si(001) planes is known to proceed faster than on Si(110),22 allowing a degree of planarization of the surface (i.e., hole filling), and the possibility to produce a flat mesh of Si(001) and Si(110) regions, Figure 1e. Figure 2a shows an atomic-force microscope (AFM) image of an etchant access hole in a 190 nm thick Si(110) NM that has been released and bonded to a Si(001) substrate, where the Si(001) host is exposed through the access hole. A comparison of the AFM scans before and after CVD overgrowth to fill the holes (Figures 2a and 2b respectively) shows that the surface has become more planar. Calibration of the differential growth rate at our growth temperature of 580˚C is obtained by measuring the step height across the boundary of the two orientations after overgrowth and, using XRD, to obtain the thickness of the Si(110)NM before and after


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overgrowth. Figure 2c shows XRD (θ/2θ) line scans of the (220) reflection from a region of the (110) NM taken before and after overgrowth. Fitting of the separation of thickness fringes for each scan yields a total growth thickness on the (110) NM of 35 nm, which, when combined with the 20nm step height after overgrowth and the original step height of 190nm, gives a Si(001) to Si(110) growth rate ratio of ~6:1 at 580°C. The combination of the original transferred membrane thickness and this growth rate anisotropy determine the overall thickness of the hybrid-orientation structure. With state-of-the-art processing, freestanding membranes less than 2nm thick are possible,21 and the upper end of the thickness range is simply dictated by the starting SOI template layer thickness (microns-thick are commercially available). If, instead of etching and rebonding a Si(110) NM on the Si(001) host of SOI(110), as above, we transfer the patterned Si(110) NM to SOI(001), we enable an easy optical first assessment of the degree of planarity of growth, as the BOX provides for color contrast (via optical interference) as a function of the Si thickness on top of it. This approach also allows a second release step, to generate a freestanding planar hybrid-orientation NM. The formation of planar structures is achieved by stopping the CVD growth at the appropriate time for a given thickness of the overlaid NM. Figure 3 shows optical-microscope images of 70 nm thick Si(110)NMs before (Figure 3a) and after (Figure 3b) overgrowth, along with schematic cross sections. In both, the Si(110) NMs were transferred to SOI(001) with a 27 nm Si template layer prior to overgrowth. The Si(110) NM is clearly visible in Figure 3a, whereas after overgrowth (Figure 3b) it becomes difficult to distinguish optically the regions of Si(001) from those of Si(110) when the heights are similar. It is also apparent from Figure 3a that NM transfer and bonding generally produces a smooth surface with negligible bubbles or trapped particulates, at least on the scale of several mm. After bonding, the NM must then survive pre-growth chemical


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cleaning, which consists of boiling-acid treatments, and an in-situ flash heating to 900°C prior to CVD growth at 580°C. Figure 3b indicates that the membrane has survived these treatments without damage, pointing to the robustness of the bond. Both the planarity of the growth across the two orientations and the RMS roughness of each region after overgrowth were assessed with AFM. Step height measurements across the boundary between (001) and (110) regions shown in Figure 3c give a height difference of only 0.5 nm between the two orientations for a 70nm (110)NM after overgrowth (timed to produce a flat growth front, using the growth ratio determined earlier). The RMS roughness is 0.4 nm on both the membrane and the overgrown (001) regions in Figure 3c, which is comparable to that of the original SOI(001) template layer. The line of defects that protrude from the interface here (note also, that the AFM tip tracks poorly over this region, which obscures the morphology) is not always observed, and indeed it is possible to obtain a very smooth transition between the two regions, as will be demonstrated later Although we exclusively use wet transfer of NMs here, dry printing and stamping are established techniques that would also be appropriate for scaling to larger dimensions and transfer to flexible polymer hosts.5,23,24,25 With dry printing, the processing steps shown in Figure 1 remain the same, except that after HF removal of the BOX, the membrane is allowed to settle on its handle substrate, where it can be adhered to a flexible polymer and printed to other substrates. This printing method could be employed to transfer the entire structure shown in Figure 2c directly to flexible polymer hosts, as is commonly done with SiNMs for use in flexible electronics and non-planar surfaces.24 Indeed, dry printing can allow complex integration in many other platforms on either the device level or wafer scale.26,27


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2.2 Structure of internal interfaces after epitaxial growth in hybrid-orientation nanomembranes Internal interfaces generated by transfer and overgrowth include a bonding interface between the Si(110) membrane and the SOI(001) host, more or less vertical interfaces occurring periodically between the side walls of the Si(110) NM and the Si(001) growing in the etchant access holes, and an epitaxial interface introduced during CVD growth of Si on the Si template layer of SOI(001). In what follows, the template layer thickness is nominally 220 nm and the thickness of the Si(110)NM is nominally 70 nm; i.e., this is a different sample from that used in Figure 3. In most cases the template layer thickness is arbitrary, it simply serves as the seed for epitaxial overgrowth. The quality of the Si(110)/(001) layers, the NM alignment, and bonding interface are critically important to ensure that the NMs survive the chemical cleaning required before introduction to the CVD growth chamber and to understand the influence of any structural defects when the NM stack is to be used for vertical charge transport.7 Therefore, we carried out a rigorous structural investigation of the samples. Figure 4a depicts an overview cross-section transmission electron microscope (TEM) image of the complete overgrown hybrid-orientation structure, showing all the types of interfaces present in the overgrown sample and allowing us to determine the various layer thickness. We confirm the initial template thickness, 210±2 nm, and determine the grown-Si thickness on the Si(001) regions to be 84±2 nm. We can also identify the three important interfaces, which we will discuss later in detail: 1) the horizontal bond interface


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between the Si(001) template and Si(110) NM, 2) the horizontal Si (001) epitaxial growth interface and 3) the vertical growth interface between Si (110) and Si (001). To verify the local information obtained by TEM, the combined Si(110) NM bonded to a SOI(001) host overgrown with Si (example shown in Figure 3b) was investigated on the macroscopic scale with X-ray reflectivity (XRR) as depicted in Figure 4b. The XRR intensity curve shows clear oscillations as a function of the diffraction vector q [q=(4π/λ)sin(θ)]; where λ is the wavelength and θ the diffraction angle in the Bragg-Brentano coplanar geometry. As the XRR signal arises from an electron density (in our case chemical) contrast and does not carry any crystal lattice information, the measurement cannot distinguish between the Si(110) membrane and the Si(001) overgrowth areas. As our TEM images indicate, as expected, a perfect epitaxial interface between the Si (001) template and the Si (001) overgrowth regimes (demonstrated in Figure 4a as well as discussed in relation to Figures 6a and 7b), we can ascribe the rapid oscillations to the interface between the bonded Si(110) NM and the Si(001) template layer of SOI(001). Using the software Motofit,28 we fitted the XRR signal using the layer model shown in the inset of Figure 4b. From the fit, we retrieve an average thickness of the interface, which we assume is SiO2, of 0.6 nm, which agrees well with TEM image shown in the inset in Figure 4b. The total thicknesses of the (110)NM (this includes the original (110) membrane and overgrowth) and the (001) template layer (again, including the original template layer and overgrowth) were 89 nm and 269 ± 40 nm, respectively. The uncertainty in the overgrown (001) template layer thickness arises primarily from the inhomogeneous thickness of the layer, and can be ascribed to the fact that XRR will not see the Si(001) template as distinct from the Si(001) overgrowth region, but rather as one Si(001) layer, as the growth interface is perfect. This gives rise to the observation that in the XRR layer model, a good fit of the XRR curve is obtained with


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assumed Si template layer thicknesses of 215 nm to 290 nm; where the minimum denotes using only the initial Si(001) template layer, and the maximum represents an upper bound on the total layer thickness. The upper thickness bound of the template layer and overgrowth is consistent with the layer thickness determined locally from the TEM images. The intensity oscillations continue to high q values, confirming that all interfaces inside the sample are smooth and have low roughness on the macroscopic scale. The inset in Figure 4b shows a high-resolution transmission electron microscope (HRTEM) image of a Si(110)/(001) bond interface. A very thin oxide layer (0.6 nm) is present at the interface. Lattice fringes are evident in the Si(110) NM, indicating that the bonding and subsequent chemical and thermal treatments leave the crystalline structure intact. The quality of the bonded interface is clearly free from fractures, debonding, or appreciable roughness on a local scale. To analyze the in-plane lattice alignment of the Si(110) membranes with the Si(001) template layer as well as to evaluate strain and crystal quality, we carried out grazing-incidence (GI) XRD. In Figures 5a and 5b a rocking curve (θ scan) and a 2θ/θ diffraction curve around the in-plane (220) reflection are shown. The zero value for the q-vector of θ scans has been chosen to be the peak position of the Si(110) membrane. In addition to the two peaks that we ascribe to the Si(110) NM and the underlying Si(001) template, we observe a diffuse background below the peak for the NM in the θ scan (Figure 5c). The diffuse background can also be observed in the 2θ/θ scan taken at the position of the NM peak (Figure 5b). To extract the position of the peaks as well as their widths, Gaussian curves were fitted to them. The contribution of each peak is plotted separately in Figures 5c and 5d; the sum of the fits is plotted against the measured data (Figures 5a and 5b) as a solid green curve. From the positions of the membrane and the template


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peak we calculate a misalignment of the crystal lattices of the Si (110) membrane and Si(001) template of 0.15° along the in-plane direction. Using the full width at half-maximum (FWHM) of the Gaussian fits to the two diffraction curves, we can extract quantitative information, such as the size of the crystal areas that give rise to the observed peaks and the background. The widths of the membrane and template peaks are of the order of the coherence length of the x-ray beam (a few micrometers). Therefore, we cannot put an upper limit on the crystal areas and conclude that the crystal structures of the template as well as the NM stay intact and are not damaged during transfer, cleaning, or overgrowth. Furthermore, we determine the width of the area giving rise to the diffuse background observed in Figures 5a and 5b to be 50-90nm. We assume the origin of this background is the vertical growth interface between the Si(110) NM and the neighboring Si(001) regions (i.e. where growth on the sidewalls of the membrane meets the growth from the template layer, as indicated in the TEM image shown in Figure 4a). Comparing the FWHMs in the θ scan and 2θ/θ scan, we conclude that the NM is strained relative to the template; as the strain is not trending towards either tensile or compressive, we ascribe it to small irregular bonding regimes between template and NM. To understand the implications of the XRD results, high-resolution TEM (HRTEM) has been carried out in these interface regions. The structure of the vertical growth interface is shown in Figures 6 and 7. In Figure 6b an outline of all the interfaces is superimposed on the crosssectional TEM micrograph shown in Figure 6a to give an overview of all the internal interfaces. The images show the overgrowth on the SOI(001) template layer (above the green area), and the lesser growth thickness on top of the Si(110) NM (regime above the blue marked area). Furthermore, we observe a band of localized defects at the horizontal interface of the SOI(001)


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template layer and its epitaxial overgrowth, denoted 3 in the schematic diagram. A similar band is evident at the top face and the sidewall of the Si(110) membrane, allowing identification of the sample structure before the overgrowth. The Si(110)/(001) vertical growth interface (marked 1) is indicated with a dashed line in Figure 6b. This border has its origin at the corner of the bonded membrane and the template layer and develops towards the Si(001) region as overgrowth on the membrane proceeds in both the horizontal and vertical (i.e., out from the sidewalls of the membrane) direction. Interestingly, we observe at the growth front near this vertical boundary two mounds of material on either side of the two growth regimes. These mounds are more evident in the overview image of Figure 4a. They indicate that the crystal growth near the interface is slower, with possibly some material transport away from this vertical interface. The observation suggests surface energy driven diffusion away from the vertical interface may be occurring, akin to grain boundary grooving during annealing of twinned crystals.29 To investigate and characterize this vertical interface between Si growing on the Si(110) NM and the Si(001) template crystal on the atomic scale, a HRTEM image was obtained exactly at the corner, where all orientations meet. It is shown in Figure 6c. For all regions lattice fringes can be observed, indicating high crystalline quality throughout. The interface between the Si (001) template and the overgrown Si(001) crystal cannot be identified even at this magnification, as the lattice fringes of the crystal run undisturbed over the growth interface, indicating perfect epitaxy. Furthermore, there exists no amorphous region at the interface between the Si(110) growth and the Si(001) growth.


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To quantify the vertical interface between the Si(110) and the Si(100) crystals as well as to infer the origin of the diffuse XRD background, we imaged the whole interface by HRTEM. Figure 6d shows the vertical growth interface in an overview HRTEM image. Notably, the vertical growth interface is narrow towards the surface (top) of the sample (as was shown in Figure 6c), but broadens closer to the template layer surface where the growth initiated (dark area in the middle of the image). This dark regime has roughly the same height as the original Si(110) bonded membrane. The white oval line in the image highlights the border between the two crystal orientations. We investigated the region in more detail by reorienting the crystal in a direction where lattice fringes can be observed in the Si(110) membrane as well as in the Si(001) region. Under these conditions two HRTEM images of the region, marked as the rectangles 1 and 2 in Figure 6d, were obtained. They are depicted in Figures 7a and 7b, respectively. Figure 7a shows the upper part of the interface. The interface appears smooth without any distortion in this region and the two orientations of grown Si, recognizable by their lattice fringes, border directly, with no additional region in between. Because this upper part of the interface is direct and free of defects, the observed diffuse background in XRD should not arise from this interface region. As discussed for Figure 6d, the lower region of the interface near to the border with the growth template shows an additional contrast. Figure 7b shows this area. The HRTEM allows us to identify the Si(110) and the Si(001), as well as a third region, by the lattice fringes. The border region is of high crystal quality as indicated by the observation of the lattice fringes, but the fringe spacing is different from the ones seen in the inner part of the Si(110) orientation and the Si(001) areas. For better visualization, a magnified area of the interface is shown in the inset of Figure 7b; see also Figure 6c for the point, where the growth interface intersects with the Si(001)


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template. This disturbed region has a size of ca. 10 - 15 nm in width and 50 - 60 nm in height. Taking the XRD results into account, which indicate an additional crystalline area with this size (giving rise to the background), we can conclude that the disturbed regime runs along the whole lower region of the vertical growth interface. Furthermore, the observed background peak in the XRD is slightly shifted from the membrane position, confirming our interpretation that the border regime observed in TEM, with a slightly different lattice parameter, is the cause of this shift. The observed shift is ascribed to a small tensile strain, which would explain the different impression of the area in the TEM, especially in the lower regime inside the area of the NM (marked by a white border in Figure 7b). We ascribe the origin of this interface disturbance to an alignment problem during overgrowth between the Si(110) membrane and the Si(001) overgrowth areas. Because of the observed small misalignment (0.15° from the XRD) and the lithographically defined border of the membrane, which will not be atomically flat, a rather small area arises in which the crystal is slightly disturbed, but still close to the (110) Si structure. By further continuation of the growth, the misalignment and roughness has been flatted by the growing crystal and a flat and therefore energetically more favorable crystal border is observed. We suggest that if the original misalignment could have been avoided, this third crystalline region would be absent. From these images, one can note the potential for lateral charge carrier transport in this structure. In a typical field effect transistor device, with a top gate, a conducting channel is localized near the surface of the semiconductor. What we have shown above suggests that an abrupt interface forms between the high-quality (110) and (001) orientations of Si, and that interfacial defects could therefore be avoided in transport across these interfaces.


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The fabrication of a dual-orientation material in the same plane using membrane transfer and overgrowth has several advantages over other techniques. Amorphization/template recrystallization (ATR) involves changing the orientation of regions of a Si(110) layer that is bonded to a Si(001) wafer by amorphizing those regions with ion implantation and then using the underlying Si(001) substrate as a seed to recrystallize the amorphized regions to Si(001) with a high-temperature annealing step.30,31 In this method, there is some concern over the crystalline quality of the Si(001) regions that have undergone ATR to change their orientation from (110). The ion implantation process typically leaves end-of-range defects at the boundary between the Si substrate and the amorphized region, which produce a band of dislocation loops at the boundary, and can generate threading dislocations in the recrystallized Si(001) regions. 30 With NM transfer and overgrowth, there is little concern for defect generation in the active-device area during the overgrowth, as we are simply employing epitaxial growth. Also, there is no need for either masking or excess material removal with a chemical-mechanical polishing step. Finally, the high temperatures required for techniques such as amorphization and templated recrystallization30 are avoided here. In situations where a strong growth rate anisotropy is absent, or where regions of entirely different materials in the surface layer are desired, patterned selective epitaxy could be an alternative approach, although an additional masking and etching/liftoff step is required.32 Masking protects the top of the overlaid membrane, with epitaxy occurring on the lower NM.

2.3 Charge transport measurements


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We extracted the mobility of electrons and holes on both the (001) and (110) orientations of these hybrid-orientation NMs (see the schematic diagram in Figure 3b) after doping different samples with phosphorus (n-type) and boron (p-type), and fabricating Hall bars on both the (001) and (110) regions in each. An example of a fabricated Hall bar is shown in Figure 8a. The dopant concentrations and room temperature mobilities extracted from Hall effect measurements are shown in Figure 8b. The dopant concentration is an order of magnitude lower for phosphorus compared to boron (see Methods), which prohibits direct mobility comparison between the two doping types. We can, however, investigate the mobility differences on the (001) and (110) regions for both electrons and holes. Figure 8b shows that the hole mobility increases by 20% when the Hall bars are fabricated on (110) compared to (001) regions, and the electron mobility shows a 100% increase when selecting (001) over (110) regions. These mobility values are consistent with those of equivalent highly doped bulk Si.33 Although the mobility measurements are not optimized, they clearly show that the hybrid-orientation NMs follow the expected trend of higher hole mobility on Si(110) and higher electron mobility on Si(001).17 The high structural quality of the vertical interfaces between the two different orientations presents an interesting platform for investigating charge transport across these interfaces and effects that may arise due to abrupt mobility changes as carriers traverse periodic orientation changes. In terms of charge scattering, we expect that these vertical grain boundaries will act similar to grain boundaries in other polycrystalline, textured thin films.34,35 2.4 Mixed-orientation, mixed-strain nanomembranes The work discussed so far involves only a mixing of orientations to create a hybridorientation semiconductor nanomembrane. However, many even more significant opportunities exist to combine NMs with different features to create novel properties. One such example is the


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incorporation of a strained semiconductor in the quilt. It is well known that strain influences carrier transport and optical properties. In the context of orientation dependent electronic transport, imparting strain to Si(110) offers a potential additional improvement in performance. Traditionally, achieving high levels of strain in Si(110) is problematic, because of the order-ofmagnitude higher density of threading dislocations in strain graded SiGe(110), relative to those in SiGe(001).19 We incorporate strain into Si(110) via use of a Si(110)/SiGe(110)/Si(110) trilayer in place of the original Si(110) NM, where the elastic strain sharing among the layers provides the strain without dislocation formation. The fabrication of tensilely strained Si using this approach has been reported generally,15,14 and specifically for Si(110).36,20 Figure 9 shows the procedure to incorporate the trilayer NM in a mixed-orientation sheet. A (110) oriented trilayer 8nmSi/39nmSi0.84Ge0.16/9nmSi NM was grown with molecular beam epitaxy (MBE). MBE is used instead of CVD because it allows growth at the lower temperature (450°C) required to maintain a sufficiently smooth growth front for the SiGe(110) orientation. 36 The SiGe layer is grown to a thickness below its kinetic critical thickness for misfit dislocation formation20,37 on top of the Si(110) template layer of SOI(110), followed by a Si(110) capping layer. At this stage only the SiGe(110) layer is under strain because the buried oxide layer serves to pin the Si(110) template layer at its bulk lattice constant. SiGe has a larger lattice constant than Si, and so is in a state of in-plane compression, as indicated by the arrows in the schematic diagram in Figure 9a. Figure 9a also shows a θ/2θ XRD line scan, in the vicinity of the (-2 2 0) reflection, of this as-grown heterostructure, which was used to confirm the thickness and composition of the layers. The scan shows a prominent SiGe diffraction peak. A clear Si peak is absent because the Si layers are so thin. The presence of distinct interference fringes, however, indicates smooth interfaces and coherency between the layers.


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The trilayer was patterned with an array of square holes and released from its handle wafer by removing the BOX, at which point the system is free to share strain by transferring some of the compressive strain from the SiGe alloy layer to create tensile strain in the Si layers. The trilayer was then bonded in place on its original Si(001) handle wafer. Figure 9b shows a schematic cross-section of a relaxed trilayer and an optical micrograph of the trilayer after release and bonding to the Si(001) host. The rippling that is visible in the image arises from the lateral expansion of the membrane that results from the strain sharing process. The narrow regions in the pattern that defines the holes (see, e.g., Figure 3a) act as stress concentrators and buckle up in a release-and-bond-in-place procedure as shown here, creating the periodic pattern of ripples.38 NMs that are transferred to foreign hosts generally do not show such rippling, as the strain can be fully shared and the structure smoothed out during the aqueous transfer.39 Figure 9c (top) shows a schematic cross-section of the bonded trilayer NM after it has been overgrown with CVD to produce a “quilt” of strained Si(110) and Si(001). The accompanying optical micrograph of the same NM after overgrowth demonstrates the ability of NMs to survive required post release chemical and thermal treatments. An AFM scan across the boundary of the two orientations (in the plane of the sample surface) is also shown in Figure 9c, and reveals that the interface contains very few defects, at least on the scale of the AFM image. The overgrowth resulted in an excellent degree of planarization (less than 0.2nm height difference). An XRD /2 scan along the (220) reflection taken after release and overgrowth is also shown in Figure 9c. It is not possible to determine the strain state of the overgrown NM by direct comparison of this scan with the similar XRD scan, shown in Figure 9a, of the as-grown but unreleased trilayer, as the subsequent overgrowth, in addition to filling in the holes that expose the Si(001) regions, results in a small increase in the thickness of the (110) membrane, producing


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more closely spaced thickness fringes than the as-released and transferred (but prior to overgrowth) counterpart. We are able to obtain the strain state in a different way, as a small portion of the original as-grown structure was left intact (i.e., not released), and subjected to the same CVD growth as the released and rebounded sections. In this region the trilayer remained in the as-grown strain state but increased its thickness in the same way as the released and rebounded parts of the trilayer NM. This scan, also presented in the XRD plot of Figure 9c, allows comparison of the strain state and crystalline quality between the as-grown (relaxed Si layers) and released (tensilely strained Si layers) trilayer, after overgrowth. The peaks in the XRD scans about the 220 reflection are associated with the out-of-plane lattice constants. Consequently, a peak shift to higher Bragg angle corresponds, via Bragg’s law, to a decrease in out-of-plane lattice constant, which translates to an in-plane strain via the elastic constants of the material. In an elastic system, the compressive in-plane strain lost from the alloy layer is completely transferred to tensile strain in the Si layers after the trilayer is released, and thus the in-plane strain in the Si(110) layers, εSi(110)can be determined from the out-of plane strain change in the alloy layer,  (i.e., from the alloy layer peak shift) through the relation

 , where

(1)

where cij are the elastic constants of the SiGe alloy.8


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The peak shift of 0.06±0.01° observed in Figure 9c for both the alloy peak and the thickness fringes, then corresponds to an out-of-plane strain change of 0.24±0.04% and in-plane tensile strain in the Si(110) layers of 0.48±0.08%. The final strain state of the system depends on the relative thicknesses of layers and the alloy composition of the stressor layer. A force balance model predicts the expected strain in the Si layers through the expression

(2)

where εm is the mismatch strain between the layers, t is the layer thickness, and M is the biaxial modulus of the layer. The strain determined from the XRD 2 scans agrees well with the strain of 0.41±0.02% predicted from Eqn. (2). The persistence of the thickness fringes and absence of any significant peak broadening through all stages of trilayer processing, shown in Figures. 9(ac) again indicates that the structural quality of the crystalline NMs is maintained through the entire release, bonding, cleaning, and overgrowth processes, so that indeed a high-quality strained-Si(110) material is combined in-plane with unstrained Si(001). Strained SiNMs indeed demonstrate mobility and drain current enhancements compared to their unstrained counterparts.40,41 Now it is possible to incorporate different orientations as well as strain states into a single NM material that can readily be transferred to foreign hosts, including other semiconductors, ceramics, and flexible soft materials.


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3. Conclusions

Patterned-silicon-nanomembrane transfer and subsequent overgrowth using the faster growth rate of Si(001), compared to Si(110), in CVD has been used to produce hybridorientation single-crystal nanomembranes consisting of a quilt of (110) and (001) regions. This approach, coupled with the ability to incorporate defect-free strained Si(110) into the architecture, provides a unique new class of material. Nanomembrane assembly, transfer, and overgrowth is generally applicable to other systems, and can be harnessed for use on a variety of different hosts. There is certainly the opportunity to combine different single-crystal materials, such as III-V semiconductors with Si, to produce materials with vastly different electrical and optical properties on a single substrate. We expect that there may be considerable potential, particularly in applications such as flexible electronics, where the direct growth of even singlecrystal Si, let alone more complex materials, is not possible. Laterally varying composition, structure, or strain in the form of a very thin sheet provide significant opportunities for fabrication of devices with unique capabilities. Significant among them will be in lateral charge transport. We have demonstrated a structurally contiguous membrane with well defined internal interfaces that should enable charge carriers to cross boundaries between different, but well controlled, crystal orientations, compositions, and strain states. Both local (TEM) and global (XRR) studies confirm that all interfaces between different crystal regions had very low disorder. HRTEM indicates well defined lattice fringes in all


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regions. In contrast to bulk-bonded interfaces, these growth interfaces do not possess the defects and amorphous regions typically introduced through bonding of dissimilar materials. Furthermore, the observation of an abrupt (vertical) interface between the high-quality (110) and (001) crystal orientations provides the potential for lateral charge carrier transport across contiguous lateral interfaces of different materials, pointing to significant implications for the development of combined electronic, photonic, thermoelectric and novel tunneling devices.

4. Methods Si(110)NM preparation: We start with SOI(110), with a (110) oriented Si template layer having a thickness of 70190 nm, with the final template thickness determined by a series of thermal oxidation and sacrificial oxide stripping in buffered oxide etchant (BOE). The buried oxide layer (BOX) is 145 nm thick. The handle wafer is Si(001). The template layer is patterned with standard photolithography and reactive ion etching (RIE) to define holes, and the BOX is selectively etched in 49% HF for 30-90 min, with the etch time depending on the size and spacing of the holes. After release, the membrane undergoes a 5 min rinse in distilled deionized water (DI) before wet transfer to a new Si(001) or SOI(001) host. Bonding to the new host consists of ramping the temperature from room temperature to 100 C over 30 min, then to 500 C over the following 30 min. Si overgrowth:


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Chemical vapor deposition (CVD) is used to grow Si with silane precursor gas (10% SiH4 with balance H2) at a flow rate of 45 standard cubic centimeters per minute (SCCM) yielding a total pressure of 26 mTorr, at a substrate temperature of 580°C. If a SiGe stressor layer is incorporated in the structure (for example in Figure 8a), solid source molecular beam epitaxy (MBE) is used to grow the alloy layer and a Si cap layer (producing a tri-layer). A growth temperature of 450˚C is required to minimize surface roughness on the Si(110) plane36 while independently controlling the growth rate (0.3-0.5Å/s).

Electrical measurements: Highly doped regions of Si(001) and Si(110) were formed by spin-on doping, followed by dopant diffusion and activation with rapid thermal annealing at 800°C in a nitrogen environment for 5 min and 15 min for boron (B-150, Honeywell) and phosphorus (P-8545, Honeywell) respectively. Samples were patterned with Hall bars and the doping concentration, ns, was determined from Hall effect measurements. For p-type samples (boron) ns = 1.7-1.8 x 1014 cm-2 and for n-type samples (phosphorus) ns = 2 x 1013 cm-2. We were not able to obtain phosphorus doping concentrations as high as boron even after longer annealing times. We expect that the volume dopant concentration is of order 1019 cm-3 for boron (and an order of magnitude lower for phosphorus) although we do not know the exact value because we do not know the dopant profile in the spin-on doped layers (which are 20-30 nm deep after diffusion and activation with RTA).


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Characterization: Atomic force microscopy (AFM) (Digital Instruments NanoscopeIIIa) was performed in intermittent-contact mode. X-ray diffraction (XRD, Phillips PanalyticalXPert PRO): Highresolution θ/2θ triple-axis line scans around the (220) reflection were used to characterize layer thicknesses, Ge composition, and out of plane strain in the strained hybrid NM structures. Transmission electron microscopy (TEM) was performed with a Tecnai F30 with 300kV at the IFW Dresden. Electron diffraction patterns were calculated using Eldisca.42 Sample crosssections were prepared for TEM using focused ion beam (FIB) milling. Preparation was conducted in an nVison (Zeiss) at the IFW Dresden following well established recipes from the vendor. After an initial electron and ion deposition of a protection layer on top of the region of interest, a lamella was defined with a 30 kV Ga ion beam. The lamella is cut free and attached to a TEM sample holder and thereafter systematically thinned using a low-energy and low-current Ga ion beam until transparency for electrons is achieved. X-ray reflectivity measurements (as well as gracing incident x-ray diffraction (GI-XRD)) were carried out at the beamline XRD2 of the Brazilian Synchrotron Light Source (LNLSCampinas) at 10.2 keV using the 4+2 circle diffractometer.

FIGURES


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Figure 1. Hybrid-orientation (110)/(001) sheet fabrication via NM transfer and overgrowth: (a) SOI(110), (b) Lithography and reactive ion etching of the Si(110) define a pattern that serves to provide etchant access holes and regions where later Si can regrow. (c) The buried oxide is selectively removed in hydrofluoric acid (HF), generating a free-floating membrane. (d) The Si(110)NM is bonded to a Si(001) substrate. (e) After CVD growth of Si over the structure, the “holes” are filled in, producing a planar mesh of Si(001)/(110).

Figure 2. CVD growth over a hybrid-orientation structure consisting of a nominally 190nm thick Si(110) membrane released and bonded to a bulk Si(001) host substrate: (a) An AFM image of the structure before overgrowth; the etchant access hole, 190nm deep, is the dark region. (b) The


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same etchant access hole after Si overgrowth with CVD. Color contrast denotes height in the AFM images. The height difference is 20nm. (c) XRD θ/2θ scans of a region of the bonded (110) NM before and after overgrowth.

Figure 3. Images of a 70 nm Si(110) NM transferred to SOI(001) with a 27 nm Si template layer. (a) Optical micrograph after transfer and bonding at 500˚C, and schematic cross-section. (b) Optical micrograph after transfer and bonding, followed by overgrowth with CVD, with schematic cross-section. (c) AFM scan across the boundary of the Si(001) and (110) regions after overgrowth. The crystal orientations are indicated on the images.


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Figure 4. TEM and XRR measurements on hybrid-orientation structure. Prior to overgrowth the structure consists of a 70 nm Si(110) membrane bonded to SOI(001) with a 220 nm template layer and 3 µm BOX. (a) Overview TEM image near the Si(001)/Si(110) interface to illustrate the thickness of the template layer and the overgrowth layer, and to determine the total layer thickness. The dark vertical stripes arise from typical strain and thickness variations in the thin TEM lamella, a consequence of TEM sample preparation. (b) XRR intensity vs scattering vector together with a fitted curve from which we deduce the parameters for the model provided in the inset. We provide the layer thickness in nm and the roughness in Ångstrom. The low roughness values evidence that interfaces remain sharp when the membrane is released. The cross-sectional HRTEM micrograph inset depicts the interface between the Si(110) membrane and the SOI(001) template layer after bonding and overgrowth.


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Figure 5. (a) Rocking curve around the (220) in-plane reflection showing the peak of the Si template as well as the bonded Si(110) NM. The green solid curve is the sum of the three curves shown in c) to fit the data. (b) 2θ/θ diffraction curve around the in-plane (220) reflection. The membrane peak sits on a diffuse background. Again, the sum of the two curves in (d) is plotted as the green curve.


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Figure 6. Cross-sectional TEM micrographs showing the vertical growth interface between Si(110) and Si(001). (a) TEM micrograph. (b) The same image from a) with an outline of the structure overlaid, showing (1) Si(110)/(001) growth interface, (2) Si(110)/(001) bonded interface and (3) Si(001)/(001) epitaxial interface. The shaded regions depict the structure as it was before CVD overgrowth (blue for the bonded Si(110) NM, green for the Si template). The black box indicates the area from which image c) was obtained. (c) HRTEM image showing the corner of the template, the bonded Si(110) NM, and the overgrown epitaxial crystal. The lattice fringes show that the border is sharp and all regions are of high crystal quality. The epitaxial interface is not observable in this scale, signifying again the high quality of the grown crystal. (d) HRTEM image the entire Si(110)/(001) vertical growth interface. The white circle marks the


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area where a disruption in the otherwise perfect crystal is observed. Here the Si (110) and Si (001) crystals connect and build a grain boundary.

Figure 7. HRTEM images of the growth interface obtained along a zone axis that allows visualization of both (110) and (001) oriented crystal lattices. (a) Upper part (rectangle marked 1 in Figure 6d). The inset shows the different lattices on both sides of the interface. (b) Lower part (rectangle marked 2 in Figure 6d). Note the extended distortion of the crystal lattice on the left side of the growth border. The inset shows a lattice image of the border between the two growth orientations. The nearly horizontal white line marks the bond interface of the Si(110) NM with


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the Si(001) template, and for reference, the box over this region represents the area depicted in Figure 6c.

Figure 8. Electrical measurements on a hybrid-orientation membrane. (a) Scanning electron microscope (SEM) image, superimposed on a (110)/(001) quilt, of a Hall bar fabricated on one square of the quilt. A Hall bar without contacting wires is shown on an adjacent square. (b) Dopant concentration and carrier mobilities for Si(001) and Si(110) regions of a hybridorientation NM.


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Figure 9. Hybrid-orientation structure incorporating strained Si at various stages of fabrication. The arrows in the schematic diagrams depict the strain state, with inward pointing and outward pointing corresponding to compression and tension respectively. (a) A sketch of the compositions and thicknesses of the layers of the as-grown (before release) heterostructure (top), as determined from the XRD line scan around the 220 reflection (bottom). (b) Optical micrograph looking down on the trilayer NM after release from the handle wafer, and subsequent bonding to a Si(001) host substrate. The sketch shows the strain sharing trilayer and one of the holes through which material deposited on the substrate will grow. (c) After overgrowth with CVD. The top-view optical micrograph (left) includes regions of Si(001) and strained Si(110), some of which are labeled. The AFM image (right), taken from a scan of the overgrown membrane, reveals a height difference of

Silicon Nanomembranes with Hybrid Crystal Orientations and Strain

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