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## **Self-Organization of Mesoscopically-Ordered Parallel Rare-Eath Silicide Nanowire Arrays on Si(110)-16×2 Surface**

Ie-Hong Hong *Department of Electrophysics, National Chiayi University, Chiayi, Taiwan*

## **1. Introduction**

198 Nanofabrication

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*Materials Science*, 2009, 206, 2155-2160.

One-dimensional nanowires have attracted significant attention due to their exotic physical properties and potential applications in nanoelectronic, nanophotonic, molecular-electronic, and magnetoelectronic devices (Agarawal, 2008; Gambardella *et al.*, 2002; Melosh *et al.*, 2003; Segovia *et al.*, 1999; Snijders, & Weitering, 2010; Yeom *et al.*, 2005; Zeng *et al.*, 2008). In particular, rare-earth metal silicide nanowires are excellent candidates as low-resistance interconnections or as fully-silicided nanoelectrodes for attaching electrically active nanostructures within a nanodevice because of their high conductivity and extremely low Schottky barrier height on *n*-type silicon and their compatibility with current silicon-based integrated circuit technology. As the continuing miniaturization encounters the lithographical limits, self-assembly (Shchukin & Bimberg, 1999) is being regarded as an effective bottom-up approach for growing one-dimensional nanowires because of its potential advantages of low cost, simplicity and high throughput. Moreover, it is believed that the basic building blocks for future complementary metal-oxide-semiconductor (CMOS) architectures will feature the integration of nanostructures based on self-assembly (Lu & Lieber, 2007; Zhirnov & Herr, 2001). In this regard, the self-assembly of epitaxial rare-earth silicide nanowires on silicon surfaces is the subject of numerous studies. It has been shown that Gd, Dy, Er, Y, and Ho can self-assemble into epitaxial silicide nanowires on a planar Si(001)-21 surface with large aspect ratios (length > 1 μm, width < 10 nm) by utilizing the anisotropic lattice mismatch between the rare-earth metal silicides and the Si substrate; these rare-earth silicide nanowires grow to arbitrary length in the direction of small (<1%) mismatch and are limited in the lateral direction of large (~8%) mismatch (Chen *et al.*, 2000, 2002; Iancu *et al.*, 2009; Lee & Kim, 2003; Liu & Nogami, 2003a; Nogami *et al*., 2001; Ohbuchi & Nogami, 2006; Ye *et al.*, 2009; Zeng *et al.*, 2008; Zhang *et al.*, 2009; Zhou *et al.*, 2006).

However, due to the alternating domains of 21 and 12 reconstruction on the neighboring single-height atomic terraces of the nominally flat Si(001) surface, rare-earth silicide nanowire always grow in two orthogonal directions on the planar Si(001) substrate. These silicide nanowires formed with two perpendicular orientations will intercept each another to terminate further growth of these orthogonal nanowires. Consequently, it is difficult to obtain parallel-aligned and long nanowires for the application in nanoelectronic devices. Therefore, several studies have attempted to use a single-domain vicinal Si(001) surface as a template to grow a large array of unidirectional rare-earth silicide nanowires (Lee *et al*., 2006; Liu & Nogami, 2003b; Preinesberger, 2005 ; Yeom *et al.*, 2005). Although a massively parallel array of straightly-aligned nanowires can be produced by such an approach, the resulting nanowires are randomly-distributed, and different in size. Even an attempt to use a highly-anisotropic vicinal Si(557) surface as a nanotemplate fails to selforganize a highly-regular parallel rare-earth silicide nanowire array over mesoscopic areas of Si(557) surface (Wanke *et al.*, 2008). For practical applications of rare-earth silicide nanowires in nanoelectronic architectures, it is of paramount importance to self-organize a massively periodic array of parallel-aligned silicide nanowires with uniform distribution, single orientation and identical size over a macroscopic area on silicon surfaces.

We have recently demonstrated that the most straightforward bottom-up approach for growing a well-ordered, highly-integrated two-dimensional networks and one-dimensional parallel arrays is to use a Si(110)–162 surface as a nanopatterned template (Hong *et al.*, 2009a, 2009b, 2011a, 2011b, 2011c). Because the 162 reconstruction consists of a unique periodic up-and-down sequence of terraces with a perfect periodicity of 5 nm and an equal width of 2.5 nm (An *et al.*, 2000), this naturally-patterned nanotemplate is thus very promising for the self-ordered growth of extremely-regular parallel nanowire arrays with high nanowire density and good uniformity over mesoscopic areas. Moreover, Si(110) wafer has been recognized as an ideal substrate for next-generation CMOS nanodevices (Yang *et al.*, 2006; Mizuno *et al.*, 2005), because CMOS transistors fabricated on Si(110) wafers exhibit both higher hole mobility and ultra-dense spatial integration than ones fabricated on Si(100) wafers (Cheng *et al.*, 2006). Thus, a larger-area self-organization of highly-integrated, wellregular, parallel epitaxial nanowire arrays on Si(110) surface can provide an efficient waferscale integration for the construction of future nanowire-based CMOS nanodevices.

In this chapter, we present that a large-scale self-organization of highly-regular, parallel arrays of epitaxial rare-earth metal silicide nanowires, covering a mesoscopic area exceeding 11 μm2, can be realized through the heteroepitaxial growth of rare-earth metal silicides along the periodic terraces of a Si(110)-162 surface. Here, we chose Gd, Ce and Er metals to self-form massively highly-ordered rare-earth silicide nanowire arrays on Si(110)-162 surface, based on the following reasons: (1) The half-filled *4f* electronic state of Gd atom contains seven spin-up electrons, Gd atoms can act as a source of spin-polarized electrons injected into silicon and create spin chains via the magnetic coupling between the large magnetic moments of *4f* electrons in the adjacent atomic chaisn (McChesney *et al*.,; Žutić, 2006). Also, GdSi1.7 epilayers exhibit magnetic ordering below 44 K (Hogg *et al.*, 2002) and Gd5(Si1−xGex)4 has a giant magnetoresistance within the range from 270 K to 290 K (Pecharsky & Gschneidner, 2001). (2) CeSix (where 1.0 ≤ x ≤ 2.0) compounds have been reported to have many intriguing physical properties, such as intermediate valence, Kondo lattice, heavy-fermion superconductivity and giant magnetoresistance (Münzenberg *et al*., 2007; Shaheen & Mendoza, 1999; Smith *et al*., 2005; Yokota *et al*., 2002); however, so far no Ce-silicide nanowires can be formed by using the anisotropic lattice-mismatched growth because of the poor lattice mismatches between silicides of the light rare-earth metal Ce and Si(001) substrate along two orthogonal crystalline directions. (3) Er silicides not only possess the lowest resistance of all rare earth metal silicides, but also exhibit anomalous Hall effect below 60 K (Pierre *et al.*, 1994) and present strong magnetic dichroism due to the large magnetocrystalline anisotropy at room temperature (Castrucci *et al.*, 1995).

Our results demonstrate the fact that our developed template-directed self-organization process based on the nanotemplate of Si(110)-162 superstructure can be applicable for the spontaneous formation of massively well-regular parallel arrays of other epitaxial rare-earth metal silicide nanowires, and open up the possibility for self-ordered, large-area, parallel arrays of magnetic atomically-precise rare-earth metal silicide nanowires on Si(110)–162 surface, which could be integrated into high-density magnetoelectronic or spintronic nanodevices. Moreover, the excellent regularity of our fabricated well-ordered parallel rareearth silicide nanowire arrays is comparable to that fabricated lithographically (Melosh *et al.*, 2003), and is much better than that produced by the anisotropic lattice mismatch growth. This advantage allows us to further carry out detailed investigations of electronic, chemical and magnetic properties of such massive self-ordered parallel rare-earth silicide nanowire arrays by means of photoemission spectroscopy (Segovia *et al.*, 1999; Yeom *et al.*, 2005) and magnetic circular dichroism (Gambardella *et al.*, 2002).

## **2. Experimental methods**

200 Nanofabrication

template to grow a large array of unidirectional rare-earth silicide nanowires (Lee *et al*., 2006; Liu & Nogami, 2003b; Preinesberger, 2005 ; Yeom *et al.*, 2005). Although a massively parallel array of straightly-aligned nanowires can be produced by such an approach, the resulting nanowires are randomly-distributed, and different in size. Even an attempt to use a highly-anisotropic vicinal Si(557) surface as a nanotemplate fails to selforganize a highly-regular parallel rare-earth silicide nanowire array over mesoscopic areas of Si(557) surface (Wanke *et al.*, 2008). For practical applications of rare-earth silicide nanowires in nanoelectronic architectures, it is of paramount importance to self-organize a massively periodic array of parallel-aligned silicide nanowires with uniform distribution,

We have recently demonstrated that the most straightforward bottom-up approach for growing a well-ordered, highly-integrated two-dimensional networks and one-dimensional parallel arrays is to use a Si(110)–162 surface as a nanopatterned template (Hong *et al.*, 2009a, 2009b, 2011a, 2011b, 2011c). Because the 162 reconstruction consists of a unique periodic up-and-down sequence of terraces with a perfect periodicity of 5 nm and an equal width of 2.5 nm (An *et al.*, 2000), this naturally-patterned nanotemplate is thus very promising for the self-ordered growth of extremely-regular parallel nanowire arrays with high nanowire density and good uniformity over mesoscopic areas. Moreover, Si(110) wafer has been recognized as an ideal substrate for next-generation CMOS nanodevices (Yang *et al.*, 2006; Mizuno *et al.*, 2005), because CMOS transistors fabricated on Si(110) wafers exhibit both higher hole mobility and ultra-dense spatial integration than ones fabricated on Si(100) wafers (Cheng *et al.*, 2006). Thus, a larger-area self-organization of highly-integrated, wellregular, parallel epitaxial nanowire arrays on Si(110) surface can provide an efficient wafer-

single orientation and identical size over a macroscopic area on silicon surfaces.

scale integration for the construction of future nanowire-based CMOS nanodevices.

magnetocrystalline anisotropy at room temperature (Castrucci *et al.*, 1995).

Our results demonstrate the fact that our developed template-directed self-organization process based on the nanotemplate of Si(110)-162 superstructure can be applicable for the

In this chapter, we present that a large-scale self-organization of highly-regular, parallel arrays of epitaxial rare-earth metal silicide nanowires, covering a mesoscopic area exceeding 11 μm2, can be realized through the heteroepitaxial growth of rare-earth metal silicides along the periodic terraces of a Si(110)-162 surface. Here, we chose Gd, Ce and Er metals to self-form massively highly-ordered rare-earth silicide nanowire arrays on Si(110)-162 surface, based on the following reasons: (1) The half-filled *4f* electronic state of Gd atom contains seven spin-up electrons, Gd atoms can act as a source of spin-polarized electrons injected into silicon and create spin chains via the magnetic coupling between the large magnetic moments of *4f* electrons in the adjacent atomic chaisn (McChesney *et al*.,; Žutić, 2006). Also, GdSi1.7 epilayers exhibit magnetic ordering below 44 K (Hogg *et al.*, 2002) and Gd5(Si1−xGex)4 has a giant magnetoresistance within the range from 270 K to 290 K (Pecharsky & Gschneidner, 2001). (2) CeSix (where 1.0 ≤ x ≤ 2.0) compounds have been reported to have many intriguing physical properties, such as intermediate valence, Kondo lattice, heavy-fermion superconductivity and giant magnetoresistance (Münzenberg *et al*., 2007; Shaheen & Mendoza, 1999; Smith *et al*., 2005; Yokota *et al*., 2002); however, so far no Ce-silicide nanowires can be formed by using the anisotropic lattice-mismatched growth because of the poor lattice mismatches between silicides of the light rare-earth metal Ce and Si(001) substrate along two orthogonal crystalline directions. (3) Er silicides not only possess the lowest resistance of all rare earth metal silicides, but also exhibit anomalous Hall effect below 60 K (Pierre *et al.*, 1994) and present strong magnetic dichroism due to the large The experiments were performed in a commercial, ultrahigh vacuum, variable-temperature STM system with a base pressure of 510-11 mbar. The device-quality flat Si(110) substrates (n-type, ~10 Ωcm) were used as one-dimensional nanotemplates for the self-organization of parallel rare-earth silicide nanowire arrays. Atomically clean Si(110)-16×2 surface was prepared by well-established annealing procedures (Hong *et al.*, 2009a, 2009b), and was confirmed by STM observation of the long-range 16×2 superstructure, as shown in Fig. 1. Parallel Gd-silicide nanowire arrays were produced by depositing ~6 Å of high purity (99.95%) Gd onto Si(110) surface at 750 o C at a deposition rate of ~0.1 Å/min and subsequently quenching to room temperature. Parallel Ce-silicide nanowire arrays were produced by depositing ~8 Å of high-purity (99.95%) Ce metals onto Si(110) surface at 400 o C at a deposition rate of ~0.2 Å/min and subsequently annealing at 700 o C for 20 minutes, then slowly cooling down to room temperature. Parallel Er-silicide nanowire arrays were produced by depositing ~10 Å of high-purity (99.95%) Er metals onto Si(110) surface at 650 o C at a deposition rate of ~0.2 Å/min and subsequently annealing at 750 o C for 30 minutes. Gd, Ce and Er metals were evaporated individually from a water-cooled electron-beam evaporator with an internal flux meter; their deposition coverages were determined *in situ* using a quartz-crystal thickness monitor. During the deposition of these rare-earth metals, the system pressure was maintained below 1 10-9 mbar. The STM measurements were acquired in the constant current mode with the atomic-resolution electrochemically etched nickel tips. The I-V curve was performed by recording the tunneling current (*It*) while ramping the sample bias (*Vs*) at a specific location on the nanowire; and was then numerically differentiated to obtain the dI/dV curve (i.e., STS spectrum).

## **3. Results and discussion**

#### **3.1 One-dimensional naturally-patterned nanotemplate of Si(110)-162 surface 3.1.1 Atomically clean Si(110)-162 surface**

Figure 1 represents typical STM topographic images of the atomically clean Si(110)–162 surface at different magnifications. As visibly seen in Figs. 1(a) and 1(b), the parallel-aligned up-and-down terraces of the 162 superstructure extending over a wide area exhibit a mesoscopic ordering and clearly reveal a unique highly-anisotropic character of this reconstruction. Such uniform grating-like terraces are very suitable to be used as a novel nanotemplate for the large-scale nanofabrication of well-ordered, high-density, parallelaligned nanowires covering macroscopic areas. A magnified image (50 50 nm2) of a part of Fig. 1(b) is shown in Fig. 1(c). It obviously shows that zigzag chains appear in the gratinglike up-and-down terraces. An atomic-resolution image (12 12 nm2) of Fig. 1(d), magnified from Fig. 1(c), shows a pair of Si pentagons forming zigzag chains on the upper terraces of 162 reconstruction are clearly resolved, consistent with the pervious result (An *et al.*, 2001; Stekolnikov *et al.*, 2004). Figure 2 displays the cross-sectional profile across the up-and-down terraces of 162 reconstruction along the line scans **C** in Fig. 1(c). The typical width and average height of these periodic upper terraces are 2.2 0.1 nm and 0.29 0.01 nm, respectively, and the regular periodicity is 5.0 0.1 nm. These measurements are consistent with the previous results (An *et al.*, 2001; Yamamoto *et al.*, 2000).

Fig. 1. (a)−(c) STM topographic images of the atomically clean Si(110)-162 surface taken at different magnifications: (a) 500 500 nm2 (*Vs* = +3 V, *It* = 30 pA), (b) 125 125 nm2, (c) 50 50 nm2, and (c) 15 15 nm2 (*Vs* = +1.5 V and *It* = 10 pA).

Fig. 2. Cross-sectional profile across the up-and-down terraces of 162 reconstruction along the white line indicated in Fig. 1(c).

nanotemplate for the large-scale nanofabrication of well-ordered, high-density, parallelaligned nanowires covering macroscopic areas. A magnified image (50 50 nm2) of a part of Fig. 1(b) is shown in Fig. 1(c). It obviously shows that zigzag chains appear in the gratinglike up-and-down terraces. An atomic-resolution image (12 12 nm2) of Fig. 1(d), magnified from Fig. 1(c), shows a pair of Si pentagons forming zigzag chains on the upper terraces of 162 reconstruction are clearly resolved, consistent with the pervious result (An *et al.*, 2001; Stekolnikov *et al.*, 2004). Figure 2 displays the cross-sectional profile across the up-and-down terraces of 162 reconstruction along the line scans **C** in Fig. 1(c). The typical width and average height of these periodic upper terraces are 2.2 0.1 nm and 0.29 0.01 nm, respectively, and the regular periodicity is 5.0 0.1 nm. These measurements are consistent

Fig. 1. (a)−(c) STM topographic images of the atomically clean Si(110)-162 surface taken at different magnifications: (a) 500 500 nm2 (*Vs* = +3 V, *It* = 30 pA), (b) 125 125 nm2, (c) 50

Fig. 2. Cross-sectional profile across the up-and-down terraces of 162 reconstruction along

with the previous results (An *et al.*, 2001; Yamamoto *et al.*, 2000).

50 nm2, and (c) 15 15 nm2 (*Vs* = +1.5 V and *It* = 10 pA).

the white line indicated in Fig. 1(c).

#### **3.1.2 Parallel silicon nanowire array grown naturally on Si(110) surface**

Since both the width and height of these upper terraces of Si(110)–162 surface are nearly close to those of silicon nanowires formed on Ag(110) surface (Sahaf *et al.*, 2007; Padova *et al.*, 2008), the long-range, parallel-aligned upper terraces of 162 reconstruction can be recognized as massively-parallel silicon nanowires grown naturally on Si(110) surface. The identification of the upper Si(110) terraces as silicon nanowires formed on Si(110) surface is similar to the self-formation of massive Si atomic lines on the Si-terminated -SiC(100)-32 surface (Soukiassian *et al.*, 1997). Silicon nanowires has been regarded as the most promising basic building blocks for the bottom-up assembly of integrated electronic and photonic nanodevices because they have the advantage of easy integration into the existing siliconbased semiconductor industry (Agarawal, 2008; Lu & Lieber, 2007).

### **3.2 Parallel Gd-silicide nanowire array self-organized on Si(110)-162 surface**

Figures 3(a)−3(e) show the representative STM topographic images of parallel Gd-silicide nanowire array at different magnifications. As clearly observed in Figs. 3(a)−3(e), these parallel-aligned, straight and nearly defect-free nanowires cover a mesoscopic area with a typical lateral separation beyond 1.0 µm and are most elongated along the [1 1 2 ] direction. In Fig. 3(d), each Gd-silicide nanowire consists of double bead chains separated by a bean chain. The double bead chains are higher than the bean chain by a difference of ~25 pm, as shown in the cross-sectional profile of Fig. 4. Each bead chain is composed of rounded protrusions marked by the red circles of ~1.5 nm in diameter, which are too large to be interpreted as individual atoms. Additionally, the adjacent substrate between the neighboring nanowires still shows a zigzag chain structure. All individual nanowires are atomically precise nanostructures, essentially identical to one another over the entire macroscopic area of Si(110) surface. Their width remains atomically precise along the nanowire and is typically 4.0 0.1 nm [see Fig. 4]. Therefore, these nanowires possess an extraordinarily high aspect ratio of greater than 250. Moreover, this massively parallel array shows a regular periodicity of 7.2 0.2 nm [see Fig. 4]. Such a perfect one-dimensional selforganized parallel nanowire array with good uniformity and alignment over a mesoscopic area represents a well-established long-range spatial ordering. This large-area Gd-silicide nanowire array also exhibits a high nanowire density and thus can be applied for the waferscale integration of high-density nanoelectronic devices. The growth mechanism of such a highly-regular parallel array of uniformly-spaced Gd-silicide nanowires on Si(110)-162 surface is driven by the heteroepitaxial growth of Gd silicides on these periodic grating-like upper terraces of Si(110)-162 superstructure, as explained in detail in our previous publication (Hong *et al.*, 2009a). This perfect self-organization of a large-area parallel array consisting of atomically-precise Gd-silicide nanowires with regular periodicity is observed for the first time.

Figures 3(e) and 3(f) show an enlarged part of the STM image in Fig. 3(d), recorded at the sample bias *Vs* = +2.0 V and −2.0 V, respectively. We note that both empty-state and filledstate STM images [i.e., Figs. 3(e) and 3(f)] represent the same area of the sample. In dual-bias images, the parallel, registry-aligned nanowires are well resolved. The empty-state image [Fig. 3(e)] clearly shows that the rounded protrusions of the double bead chains in Fig. 3(d) converts to pentagons; however, the pentagons in the left and right bead chains are oppositely oriented, which is the same as the pair of Si pentagons of Si(110)-162 superstructure [see Fig. 1(d)]. Moreover, the bean chain disappears and instead becomes to a trench. This result strongly suggests that the Si pentagon pair on the upper terraces of 162 reconstruction was split into two individual Si pentagons due to a strong reaction of Si and Gd upon the adsorption of Gd on the upper terraces. Consequently, a trench was formed in the middle of Gd-silicide nanowires and Gd atoms were embedded in the trench (i.e., Gd atoms were adsorbed in the bean chain of Fig. 3(d)), similar to the Si-deficient structural model (Lee & Kim, 2003). The width of Gd-silicide nanowire thus becomes to be ~4 nm from the Si terrace width of ~2.2 nm. In the filled-state image [Fig. 3(f)], each Gd-silicide nanowire appears to consist of two sawtooth rows of distinct atomic arrangements on both sides and a pair of double linear chains with an up-down buckling configuration in the middle of nanowire; moreover, the contrast of these two sawtooth rows is much lower than that of double linear chains. Notice that the dark trench between the double bead chains in Fig. 3(e) inverts to the bright double linear chains in Fig. 3(f) when the bias polarity is reversed. Considering the charge transfer from Si to Gd atoms by electronegativity, the bright double linear chains in Fig. 3(f) are primarily contributed from the electronic states localized on Gd atoms, suggesting that Gd atoms are located on the bean chain in Fig. 3(d). That is, the two sawtooth rows and the double linear atomic chains consist mainly of Si and Gd atoms, respectively. The appearance of the well-ordered internal structures in Gd-silicide nanowires strongly verifies that these straightly-aligned, long Gd-silicide nanowires grow epitaxially on Si(110) surface.

Fig. 3. (a)−(d) STM topographic images of parallel Gd-silicide nanowire array on Si(110)- 162 surface taken at different magnifications: (a) 0.8 0.8 μm2 (*Vs* = +3 V, *It* = 30 pA), (b) 250 250 nm2, (c) 110 110 nm2, and (d) 50 50 nm2. (e), (f) Dual-polarity images (25 25 nm2) of parallel Gd-silicide nanowire array acquired at +2 V and −2 V, respectively, and at *It* = 10 pA.

trench. This result strongly suggests that the Si pentagon pair on the upper terraces of 162 reconstruction was split into two individual Si pentagons due to a strong reaction of Si and Gd upon the adsorption of Gd on the upper terraces. Consequently, a trench was formed in the middle of Gd-silicide nanowires and Gd atoms were embedded in the trench (i.e., Gd atoms were adsorbed in the bean chain of Fig. 3(d)), similar to the Si-deficient structural model (Lee & Kim, 2003). The width of Gd-silicide nanowire thus becomes to be ~4 nm from the Si terrace width of ~2.2 nm. In the filled-state image [Fig. 3(f)], each Gd-silicide nanowire appears to consist of two sawtooth rows of distinct atomic arrangements on both sides and a pair of double linear chains with an up-down buckling configuration in the middle of nanowire; moreover, the contrast of these two sawtooth rows is much lower than that of double linear chains. Notice that the dark trench between the double bead chains in Fig. 3(e) inverts to the bright double linear chains in Fig. 3(f) when the bias polarity is reversed. Considering the charge transfer from Si to Gd atoms by electronegativity, the bright double linear chains in Fig. 3(f) are primarily contributed from the electronic states localized on Gd atoms, suggesting that Gd atoms are located on the bean chain in Fig. 3(d). That is, the two sawtooth rows and the double linear atomic chains consist mainly of Si and Gd atoms, respectively. The appearance of the well-ordered internal structures in Gd-silicide nanowires strongly verifies that these straightly-aligned, long Gd-silicide nanowires grow

Fig. 3. (a)−(d) STM topographic images of parallel Gd-silicide nanowire array on Si(110)- 162 surface taken at different magnifications: (a) 0.8 0.8 μm2 (*Vs* = +3 V, *It* = 30 pA), (b) 250 250 nm2, (c) 110 110 nm2, and (d) 50 50 nm2. (e), (f) Dual-polarity images (25 25 nm2) of parallel Gd-silicide nanowire array acquired at +2 V and −2 V, respectively, and at *It* = 10

epitaxially on Si(110) surface.

pA.

As seen in Figs. 3(e) and 3(f), the two bead chains of empty orbitals on both sides of nanowire are in antiphase with the double linear chains of filled orbitals in the middle of nanowire. This polarity dependence of STM images clearly reveals that each Gd-silicide nanowire really consists of a bundle of chain-like structures with a peculiar charge arrangement of alternating filled and empty orbitals. These dual-polarity STM images clearly show that these parallel-aligned Gd-silicide nanowires are well-defined and periodically arranged in both filled-state and empty-state images. The periodic intensity modulation exhibited by this parallel array in dual-bias images clearly reveals a perfect spatial-organization of the charge distribution. Such a well-regular charge modulation of this parallel Gd-silicide nanowire array could be exploited in nanoelectronic devices.

Figure 4 displays the cross-sectional profiles of the line scan **A** across the parallel Gd-silicide nanowire array in Fig. 3(d). This cross-sectional profiles show clearly that the shape distribution of the parallel Gd-silicide nanowire array is well-defined and uniformly-spaced. The unprecedented uniformity exhibited by this parallel array reflects a perfect spatialorganization characteristic of our performed one-dimensional self-assembling process. As shown in the topographic profile in Fig. 4, the nanowires are typically 4.0 0.1 nm in width and have an average periodicity of 7.2 0.2 nm; the heights observed in the empty-state image and the filled-state image are, respectively, 270 10 pm and 250 10 pm, similar to the height obtained by Lee *et al* (Lee & Kim, 2003). It is noted that the nanowire height is nearly equal to that of Si upper terraces on Si(110)-162 surface, which suggests that Gd atoms may incorporate into Si upper terraces to form one-dimensional Gd-silicide nanowires, similar to the proposed structure of Si-deficient rare-earth metal silicide nanowires (Chen *et al.*, 2002; Eames *et al.*, 2010).

Fig. 4. Cross-sectional profile **A** across the parallel Gd-silicide nanowire array along the white line indicated in Fig. 3(d)

To gain insight into the local electronic structure of Gd-silicide nanowires, STS spectra (i.e., dI/dV curves) were measured on top of the individual Gd-silicide nanowires and the adjacent substrate between neighboring nanowires, as displayed in Fig. 5. A third dI/dV curve taken on clean Si(110) surface is also shown for reference. Each dI/dV curve is laterally averaged over 20 individual spectra taken within the corresponding region. These STS spectra allow us to compare the surface conductivity and local density of states (LDOS) in the areas selected on the surface. Curve **I** measured on the self-organized parallel Gdsilicide nanowires show non-zero conductance near the zero bias, indicating that Gd-silicide nanowires have one-dimensional metallic character along their length. As clearly revealed in Curve **I**, the Gd-silicide nanowires are characterized by a set of the discrete, distinct electronic states.

Fig. 5. Averaged dI/dV curves of the parallel Gd-silicide nanowires (Curve **I**), the adjacent substrate (Curve **II**), and the clean Si(110) surface (Curve **III**). Each dI/dV curve was laterally averaged over their corresponding areas in the inserted STM image (15 10 nm2).

These sharp peaks in the STS spectrum of individual Gd-silicide nanowire are characteristic of the Van Hope singularities expected for the LDOS of one-dimensional materials (Iancu *et al.*, 2009; Zeng *et al.*, 2008). Curve **II** taken over the adjacent substrate shows semiconducting behavior with a band gap of ~0.5 eV, significantly smaller than ~1.2 eV for clean Si(110) surface [see Curves **III**]. The reduced gap opening of the substrate together with the zigzag chain structure suggests that Si(110)-162 surface is reconstructed with Gd.

#### **3.3 Parallel Ce-silicide nanowire arrays self-organized on Si(110)-162 surface**

Figure 6 shows the STM topographic images of the parallel Ce-silicide nanowire array at different magnifications. As visibly observed in Figs. 6(a)−6(e), these parallel-aligned, straight and nearly defect-free nanowires cover a mesoscopic area with a typical lateral distance exceeding 1.3 μm and are elongated along the [1 1 2 ] direction. These nanowires thus possess an extraordinarily high aspect ratio. Also, this massively parallel nanowire array shows a regular periodicity and a high integration density. Moreover, these parallel nanowires are essentially identical to one another over the entire macroscopic area of the Si(110) surface. This result clearly reveals that these parallel nanowires grew along the periodic terraces of Si(110)-162 superstructure. Such large-area parallel Ce-silicide nanowire array could be applied in the wafer-scale integration of high-density nanoelectronic architectures on Si(110) surface. As clearly seen in Fig. 6(e), each Ce-silicide nanowire consists of double zigzag chains with distinct morphologies. In the atomic-

in the areas selected on the surface. Curve **I** measured on the self-organized parallel Gdsilicide nanowires show non-zero conductance near the zero bias, indicating that Gd-silicide nanowires have one-dimensional metallic character along their length. As clearly revealed in Curve **I**, the Gd-silicide nanowires are characterized by a set of the discrete, distinct

Fig. 5. Averaged dI/dV curves of the parallel Gd-silicide nanowires (Curve **I**), the adjacent substrate (Curve **II**), and the clean Si(110) surface (Curve **III**). Each dI/dV curve was laterally averaged over their corresponding areas in the inserted STM image (15 10 nm2). These sharp peaks in the STS spectrum of individual Gd-silicide nanowire are characteristic of the Van Hope singularities expected for the LDOS of one-dimensional materials (Iancu *et al.*, 2009; Zeng *et al.*, 2008). Curve **II** taken over the adjacent substrate shows semiconducting behavior with a band gap of ~0.5 eV, significantly smaller than ~1.2 eV for clean Si(110) surface [see Curves **III**]. The reduced gap opening of the substrate together with the zigzag

chain structure suggests that Si(110)-162 surface is reconstructed with Gd.

**3.3 Parallel Ce-silicide nanowire arrays self-organized on Si(110)-162 surface**

Figure 6 shows the STM topographic images of the parallel Ce-silicide nanowire array at different magnifications. As visibly observed in Figs. 6(a)−6(e), these parallel-aligned, straight and nearly defect-free nanowires cover a mesoscopic area with a typical lateral distance exceeding 1.3 μm and are elongated along the [1 1 2 ] direction. These nanowires thus possess an extraordinarily high aspect ratio. Also, this massively parallel nanowire array shows a regular periodicity and a high integration density. Moreover, these parallel nanowires are essentially identical to one another over the entire macroscopic area of the Si(110) surface. This result clearly reveals that these parallel nanowires grew along the periodic terraces of Si(110)-162 superstructure. Such large-area parallel Ce-silicide nanowire array could be applied in the wafer-scale integration of high-density nanoelectronic architectures on Si(110) surface. As clearly seen in Fig. 6(e), each Ce-silicide nanowire consists of double zigzag chains with distinct morphologies. In the atomic-

electronic states.

resolution image of Fig. 6(f), we can visibly find that the elemental structures of the double zigzag chains in the Ce-silicide nanowire are completely different from those of up-anddown terraces on Si(110)-162 surface; moreover, the width of all individual nanowires remains atomically precise along their lengths. This perfect one-dimensional self-organized parallel array with good uniformity and alignment over a mesoscopic area of 1.3×1.3 μm2 reveals a well-established long-range spatial ordering. The spontaneous formation of such a large-area, highly-ordered parallel array consisting of atomically-precise Ce-silicide nanowires with a regular periodicity has never been reported before.

Fig. 6. A series of different magnified STM topographic images of the parallel Ce-silicide nanowire array: (a) 1.3 1.3 μm2 (*Vs* = 2.0 V, *It* = 0.1 nA); (b) 850 850 nm2; (c) 250 250 nm2; (d) 125 125 nm2; (e) 50 50 nm2 (*Vs* = 2.0 V, *It* = 20 pA); and (f) 35 35 nm2.

Figures 7(a) and 7(b) show dual-polarity STM images of an enlarged area of the parallel Cesilicide nanowire array in Fig. 6(e), recorded at the sample bias *Vs* = +1.5 V and −1.5 V. The empty-state image [Fig. 7(a)] at +1.5 V clearly shows a set of double zigzag chains on each Ce-silicide nanowire and the right zigzag chain appears much brighter than the left chain. In the filled-state image [Fig. 7(b)] at −1.5 V, the Ce-silicide nanowires consist of two linear rows of distinct atomic arrangements and the right linear row is also brighter than the left row. These dual-bias STM images evidently show that Ce-silicide nanowires are registryaligned and really consist of a bundle of double chain-like structures with different morphologies. Figure 7(c) plots the superposition of the cross-sectional profiles of both line scans **E** and **F** across the empty-state and filled-state images of parallel Ce-silicide nanowires in Figs. 7(a) and 7(b), respectively. The section profiles of both lines **E** and **F** clearly show that these parallel Ce-silicide nanowires are well defined and periodically positioned. The unprecedented regularity exhibited by this parallel array in dual-bias images also reflects a perfect spatial-ordering of the charge distribution.

As clearly revealed in Fig. 7(c), these parallel-aligned nanowires are identical to 5.0 0.2 nm in width and have an equal periodicity of 6.0 0.2 nm. Moreover, the topographic maxima of the double zigzag chains in the empty-state image and the double linear rows in the filled-state image are spatially coincident; indicating that the height difference between the right chains and left chains is due to the topographic effect, not the electronic effect. The left–right asymmetry observed in the dual-polarity images of each Ce-silicide nanowire indicates that the building blocks for Ce-silicide nanowires are asymmetrically composed. Additionally, the widths of the left and right zigzag chains are 2.0 0.2 and 3.0 0.2 nm, respectively; while the double linear rows are equal to 2.5 0.1 nm in width, similar to the width (2.2 0.1 nm) of Si terraces of the Si(110)-162 reconstruction [see Fig. 2].

Fig. 7. Dual polarity STM images (45 35 nm2) of the parallel Ce-silicide nanowires acquired at +1.5 V (a) and −1.5 V (b), respectively, and at 20 pA. (c) Cross-sectional profiles **E** and **F**  across parallel-aligned Ce-silicide nanowires of the empty- and filled-state images along the white lines indicated in (a), respectively.

These results strongly suggest that Ce atoms nucleated concurrently along the upper and lower terraces of Si(110)-162 surface to form the one-dimensional Ce-silicide nanowires consisting of double chain row structures with different heights. The Ce-silicide nanowires are characterized by a well-defined shape indicating epitaxial growth, which is promising for the basic building block of nanoelectronic devices.

It is very interesting to explore the local electronic structures of the parallel Ce-silicide nanowire array. STS spectra (i.e., dI/dV curves) were measured on the individual Cesilicide nanowires and also on the adjacent substrate between the neighboring nanowires, as displayed in Fig. 8. A third dI/dV curve taken on clean Si(110) surface is also shown for comparison. Each dI/dV curve was laterally averaged over 30 individual spectra taken within the corresponding areas indicated in the inserted STM image (20 10 nm2). Curve **I** for the left chain shows an energy gap of 0.68 eV, suggesting a semiconducting property

As clearly revealed in Fig. 7(c), these parallel-aligned nanowires are identical to 5.0 0.2 nm in width and have an equal periodicity of 6.0 0.2 nm. Moreover, the topographic maxima of the double zigzag chains in the empty-state image and the double linear rows in the filled-state image are spatially coincident; indicating that the height difference between the right chains and left chains is due to the topographic effect, not the electronic effect. The left–right asymmetry observed in the dual-polarity images of each Ce-silicide nanowire indicates that the building blocks for Ce-silicide nanowires are asymmetrically composed. Additionally, the widths of the left and right zigzag chains are 2.0 0.2 and 3.0 0.2 nm, respectively; while the double linear rows are equal to 2.5 0.1 nm in width, similar to the

Fig. 7. Dual polarity STM images (45 35 nm2) of the parallel Ce-silicide nanowires acquired at +1.5 V (a) and −1.5 V (b), respectively, and at 20 pA. (c) Cross-sectional profiles **E** and **F**  across parallel-aligned Ce-silicide nanowires of the empty- and filled-state images along the

These results strongly suggest that Ce atoms nucleated concurrently along the upper and lower terraces of Si(110)-162 surface to form the one-dimensional Ce-silicide nanowires consisting of double chain row structures with different heights. The Ce-silicide nanowires are characterized by a well-defined shape indicating epitaxial growth, which is promising

It is very interesting to explore the local electronic structures of the parallel Ce-silicide nanowire array. STS spectra (i.e., dI/dV curves) were measured on the individual Cesilicide nanowires and also on the adjacent substrate between the neighboring nanowires, as displayed in Fig. 8. A third dI/dV curve taken on clean Si(110) surface is also shown for comparison. Each dI/dV curve was laterally averaged over 30 individual spectra taken within the corresponding areas indicated in the inserted STM image (20 10 nm2). Curve **I** for the left chain shows an energy gap of 0.68 eV, suggesting a semiconducting property

white lines indicated in (a), respectively.

for the basic building block of nanoelectronic devices.

width (2.2 0.1 nm) of Si terraces of the Si(110)-162 reconstruction [see Fig. 2].

along its length. These results clearly reveal that self-organized parallel Ce-silicide nanowires are semiconducting. Curve **II** for the adjacent substrate shows a lager energy gap of 0.9 eV, smaller than the gap of ~1.2 eV obtained for the clean Si(110)-162 substrate [see Curve **III**].

Fig. 8. Averaged STS spectra of the right zigzag chain taken on the individual Ce-silicide nanowires (Curve **I**), the adjacent substrate between neighboring nanowires (Curve **II)**, and the clean Si(110) surface (Curve **III**).

The reduced gap opening of the substrate and the featureless LDOS means that the Si(110) surface was covered with a small amount of Ce. The semiconducting behavior of Ce-silicide nanowires is in contrast to the metallic character of other rare-earth silicide nanowires (Hong *et al.*, 2011a; Lee & Kim, 2003; Yeom *et al.*, 2005; Zeng *et al.*, 2008) and CeSix Kondo compounds (Malterre *et al.*, 1993; Mimura *et al.*, 2005). The suppression of LDOS near the Fermi energy (i.e., zero bias) can be ascribed to the formation of a Mott-Hubbard insulating states at these Ce-silicide nanowires (Hong *et al.*, 2011c).

## **3.4 Parallel Er-silicide nanowire arrays self-organized on Si(110)-162 surface**

Figure 9 shows a series of different magnified STM topographic images of parallel Ersilicide nanowire arrays. As clearly observed in Fig. 9, these parallel-aligned, straight andnearly defect-free nanowires cover a mesoscopic area with a typical length beyond 1.0

µm and are most elongated along the [1 1 2 ] direction. In Fig. 9(e), each Er-silicide nanowire consists of double bead chains, but unlike to the morphology of Gd-silicide nanowire with the bean chain in the middle of nanowire. Each bead chain is composed of round protrusions with a typical diameter of ~2.0 nm. Additionally, the adjacent substrate between the neighboring nanowires still shows a zigzag chain structure. All individual nanowires are atomically precise nanostructures, essentially identical to one another over the entire macroscopic area of Si(110) surface.Their width remains atomically precise along the nanowire and is typically 4.2 0.1 nm [see Fig. 10].Therefore, these nanowires possess an extraordinarily high aspect ratio of greater than 250. Moreover, this massively parallel array shows a regular periodicity of 8.0 0.2 nm [see Fig. 10]. Such a perfect onedimensional self-organized parallel nanowire array with good uniformity and alignment over a mesoscopic area represents a well-established long-range spatial ordering. This largearea Er-silicide nanowire array also exhibits a high nanowire density and thus can be applied for the wafer-scale integration of high-density nanoelectronic devices.

Fig. 9. (a)−(d) STM topographic images of parallel Er-silicide nanowire array on Si(110)-162 surface taken at different magnifications: (a) 1 1 μm2 (*Vs* = +3 V, *It* = 30 pA), (b) 500 500 nm2, (c) 200 200 nm2, (d) 110 110 nm2, (e) 60 60 nm2 (*Vs* = +3 V, *It* = 10 pA), and (f) 36 36 nm2 (*Vs* =+2 V, *It* = 10 pA).

Figure 10 displays the cross-sectional profiles of the line scan B across the parallel Er-silicide nanowire array in Fig. 9(e). This cross-sectional profiles show clearly that the shape distribution of the parallel Er-silicide nanowire array is well-defined and uniformly-spaced. The unprecedented uniformity exhibited by this parallel array reflects a perfect spatialorganization characteristic of our performed one-dimensional self-assembling process. As shown in the topographic profile in Fig. 10, the nanowires are typically 4.2 0.1 nm in width and have a representative periodicity of 8.0 0.2 nm; the average heights is 230 20 pm, similar to the height obtained by Chen *et al*. (Chen *et al.*, 2002) The growth mechanism of such a highly-regular parallel array of uniformly-spaced Er-silicide nanowires on Si(110)- 162 surface is also driven by the heteroepitaxial growth of Er silicides along the perfect grating-like upper terraces of Si(110)-162 superstructure, similar to the self-organized growth of Gd-silicide nanowires on Si(110)-162 surface.

array shows a regular periodicity of 8.0 0.2 nm [see Fig. 10]. Such a perfect onedimensional self-organized parallel nanowire array with good uniformity and alignment over a mesoscopic area represents a well-established long-range spatial ordering. This largearea Er-silicide nanowire array also exhibits a high nanowire density and thus can be

Fig. 9. (a)−(d) STM topographic images of parallel Er-silicide nanowire array on Si(110)-162 surface taken at different magnifications: (a) 1 1 μm2 (*Vs* = +3 V, *It* = 30 pA), (b) 500 500 nm2, (c) 200 200 nm2, (d) 110 110 nm2, (e) 60 60 nm2 (*Vs* = +3 V, *It* = 10 pA), and (f) 36

Figure 10 displays the cross-sectional profiles of the line scan B across the parallel Er-silicide nanowire array in Fig. 9(e). This cross-sectional profiles show clearly that the shape distribution of the parallel Er-silicide nanowire array is well-defined and uniformly-spaced. The unprecedented uniformity exhibited by this parallel array reflects a perfect spatialorganization characteristic of our performed one-dimensional self-assembling process. As shown in the topographic profile in Fig. 10, the nanowires are typically 4.2 0.1 nm in width and have a representative periodicity of 8.0 0.2 nm; the average heights is 230 20 pm, similar to the height obtained by Chen *et al*. (Chen *et al.*, 2002) The growth mechanism of such a highly-regular parallel array of uniformly-spaced Er-silicide nanowires on Si(110)- 162 surface is also driven by the heteroepitaxial growth of Er silicides along the perfect grating-like upper terraces of Si(110)-162 superstructure, similar to the self-organized

36 nm2 (*Vs* =+2 V, *It* = 10 pA).

growth of Gd-silicide nanowires on Si(110)-162 surface.

applied for the wafer-scale integration of high-density nanoelectronic devices.

Fig. 10. Cross-sectional profile **B** across the parallel Er-silicide nanowire array along the white line indicated in Fig. 9(e)

Figure 11 displays the representative STS spectra of the self-organized Er-silicide nanowires (Curve **I**) and the adjacent substrate between neighboring nanowires (Curve **II**). Each dI/dV curve is laterally averaged over 30 individual spectra taken within the corresponding region of the inserted STM image. Curves **I** shows non-zero conductance near the zero bias, indicating that Er-silicide nanowires exhibit metallic nature along their length. Curve **II** depicts metallic-like behavior with a narrower band gap of ~0.12 eV, significantly smaller than ~1.2 eV for the clean Si(110) surface [see Curves **III**]. The metallic-like character of the substrate together with the zigzag chain structure suggests that Si(110) surface is reconstructed with Er.

Fig. 11. Averaged dI/dV curves of the parallel Er-silicide nanowires (Curve **I**), the adjacent substrate (Curve **II**), and the clean Si(110) surface (Curve **III**). Each dI/dV curve was laterally averaged over their corresponding areas in the inserted STM image (25 12 nm2).

#### **3.5 Parallel transition metal silicide nanowires self-assembled on Si(110)-162 surface**  As clearly demonstrated in Figs. 3-11, the massively parallel arrays of various epitaxial rareearth metal silicide nanowires on Si(110)-162 surface always grow along a unidirectional

orientation of [1 1 2 ]with uniform distribution and identical size. However, the growth mode of rare-earth metal silicide nanowires on Si(110) surface is completely different from that of the transition-metal silicide nanowires on Si(110) surface. As shown in Fig. 12, the parallel-aligned Ni-silicide and Fe-silicide nanowires can not form well-ordered parallel nanowire arrays with a regular periodicity and identical dimensions on Si(110) surface, as reported in the previous results (He *et al*., 2004; Liang *et al*., 2006). The difference can be attributed to the fact that transition-metal silicide nanowires prefer to grow into the Si(110) substrate along the [1 1 0 ] direction via an "endotaxial" growth mechanism (He *et al*., 2004) and do not favor to nucleate along the [1 1 2 ]-oriented terraces of Si(110)-162 reconstruction.

Fig. 12. STM topographic images of parallel Ni-silicide nanowires on Si(110) surface [(a)−(c)] and parallel Fe-silicide nanowires on Si(110) surface [(d)−(f)] at different magnifications: (a) 7 7 μm2, (b) 1 1 μm2, (c) 500 500 nm2, (d) 800 800 nm2, (e) 125 125 nm2, and (f) 45 45 nm2.

### **4. Conclusion**

We have successfully developed a simple and efficient bottom-up nanofabrication for the self-organization of mesoscopically-ordered parallel arrays, consisting of uniformly-spaced identical-size epitaxial rare-earth metal silicide nanowires over a mesoscopic area on Si(110)- 16×2 surface. Three perfectly-ordered parallel arrays of atomically-precise Gd-silicide nanowires, Ce-silicide nanowires, and Er-silicide nanowires are presented to demonstrate the versatility of this novel approach. These massively parallel rare-earth silicide nanowire arrays were self-organized with widths and pitches (center-to-center distances) as small as 4~5 nm and 6~8 nm, respectively, and lengths exceeding 1 µm, which were achieved through the heteroepitaxial growth of rare-earth metal silicides along the periodic terraces of the Si(110)-162 surface. The ability to form such large-area, highly-ordered parallel arrays of epitaxial rare-earth metal silicide nanowires on the Si(110) nanotemplate represents a simple step towards the bottom-up nanofabrication of high-density, well-defined interconnections and functional nanowire-based CMOS nanodevices in a straightforward, cost-effective and high-throughput process. As they have the advantage of easy integration into the existing silicon-based integrated-circuit technology, we believe that our developed large-scale self-organization of epitaxial rare-earth metal silicide nanowire arrays on Si(110)- 16×2 surface will be an important Si-compatible nanofabrication process for highlyintegrated, well-defined parallel active nanoarchitectures with diverse functions over mesoscopic areas, which could be used to realize the future innovations in Si-based nanoelectronics.

## **5. Acknowledgements**

This work is financially supported by the National Science Council of Taiwan under Grant Nos. 97-2738-M-415-001 and 97-2112-M-415-003-MY3.

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Increasing attention has been paid to the control over surface chemical property of bare metals, semiconductors, ceramics, and polymers in the development of nanoscience and nanotechnology. The appropriate functionalization has a potential to lead to the appearance of new function. Organic monolayers, i.e., self-assembled monolayers (SAMs), are simple monomolecular system, convenient, and flexible to tailor the substrate's surface chemical and physical properties including wetting property, conductivity, and thermal, chemical, photochemical stabilities, etc. In the following molecular diagram of organic monolayers, its molecular anatomy and characteristics are highlighted as shown in Figure 1. Such monolayer structures are formed by chemisorption of organic constituents onto the specific solid substrate through liquid or gas phase. A headgroup of the organic constituent reacts to the outermost atoms on the substrate to give a chemical linkage. The chemisorbates organize epitaxially to form crystalline monomolecular film structures. As shown in the molecular

Fig. 1. Schematic diagram of organic monolayers formed on solid substrates

diagram, there is a number of headgroups which can recognize the outermost chemical groups of substrates, resulting in a chemical bond. In contrast, there is no limitation in the selection of terminal functional groups. A wide variety of functionalities or reactivities are applicable. A lateral molecular space, which is controlled by hydrophobic molecular interaction between hydrocarbon chains, allows the control of the film crystallinity and electron transferring behavior between the terminal group and substrate surface. Thus, programmed surface can be designed by careful control of monolayer attachment. Since the finding of siloxane SAM structure in 1980, a number of research efforts have been performed. In this review article, organic monolayer systems on polymer and silicon substrates are focused, and their potential applications is also demonstrated.

## **2. Chemical modification of a polymer surface**

#### **2.1 Motivation**

Recently, the deposition of inorganic materials, e.g., ceramics, metals, and semiconductors, has interestingly attracted attention in the development of mechanically-flexible device applications including OLED, flexible flat panel and wearable displays, electronic papers, biomedical tools. These technological developments take advantage of polymer's excellent properties including mechanical flexibility, lightweight, ease-of-design and coloring, lowcost and good impact resistance. In order to fabricate such flexible device applications, it is a key to control chemical properties of polymer's surface that induce the nucleus formation of inorganic crystals, and accelerates subsequent crystalline growth or film growth. Organosilane SAMs have been frequently employed to modify surface chemically properties of inorganic substrates [1-5]. The resultant surfaces play important roles in molecular recognition events. For example, some of functional groups induce the nucleus formation of inorganic crystals, and accelerates subsequent crystalline growth or film growth. This section describes the preparation of well-ordered SAM structures on polymer sheet, and then the deposition of metal and ceramics films on the SAM-covered polymer sheets through solution processes.

#### **2.2 SAM formation**

When silane molecules are treated directly on as-received polymer sheet, they cannot form a monomolecular structure. Similar can be seen even on photochemically treated polymer sheet. In other words, when silane precursors are directly treated on surface-activated polymer substrates, well-ordered SAM formation cannot effectively proceed, since the polar-functional groups formed on a polymer surface through surface modification methods using plasma or UV light are generally inhomogeneous and randomly distributed [5]. Such a heterogeneous surface cannot provide adequate support for the preparation of a wellordered SAM. As an example, Table 1 summarizes the contact angles and surface free energies of polyimide surfaces covered with NH2- and CF3-terminated SAMs, and their values are compared with those of SAM-covered silicon substrates used as standards. The formation of the SAMs is performed on photooxidized polyimide and silica-covered silicon substrates, respectively. The details of the SAM preparation method are described in elsewhere [5], but the surface energies of each sample are calculated using the following equation:

$$\gamma\_{\mathsf{L}}(\mathsf{1} + \cos \theta\_{\mathsf{S}} \mathsf{1}) = \mathsf{2} [ (\gamma\_{\mathsf{s}} \mathsf{4} \bullet \gamma\_{\mathsf{L}} \mathsf{4})^{1/2} + (\gamma\_{\mathsf{s}} \mathsf{1} \bullet \gamma\_{\mathsf{L}} \mathsf{1})^{1/2} + (\gamma\_{\mathsf{s}} \mathsf{4} \bullet \gamma\_{\mathsf{L}} \mathsf{1})^{1/2}]$$

diagram, there is a number of headgroups which can recognize the outermost chemical groups of substrates, resulting in a chemical bond. In contrast, there is no limitation in the selection of terminal functional groups. A wide variety of functionalities or reactivities are applicable. A lateral molecular space, which is controlled by hydrophobic molecular interaction between hydrocarbon chains, allows the control of the film crystallinity and electron transferring behavior between the terminal group and substrate surface. Thus, programmed surface can be designed by careful control of monolayer attachment. Since the finding of siloxane SAM structure in 1980, a number of research efforts have been performed. In this review article, organic monolayer systems on polymer and silicon

Recently, the deposition of inorganic materials, e.g., ceramics, metals, and semiconductors, has interestingly attracted attention in the development of mechanically-flexible device applications including OLED, flexible flat panel and wearable displays, electronic papers, biomedical tools. These technological developments take advantage of polymer's excellent properties including mechanical flexibility, lightweight, ease-of-design and coloring, lowcost and good impact resistance. In order to fabricate such flexible device applications, it is a key to control chemical properties of polymer's surface that induce the nucleus formation of inorganic crystals, and accelerates subsequent crystalline growth or film growth. Organosilane SAMs have been frequently employed to modify surface chemically properties of inorganic substrates [1-5]. The resultant surfaces play important roles in molecular recognition events. For example, some of functional groups induce the nucleus formation of inorganic crystals, and accelerates subsequent crystalline growth or film growth. This section describes the preparation of well-ordered SAM structures on polymer sheet, and then the deposition of metal and ceramics films on the SAM-covered polymer sheets

When silane molecules are treated directly on as-received polymer sheet, they cannot form a monomolecular structure. Similar can be seen even on photochemically treated polymer sheet. In other words, when silane precursors are directly treated on surface-activated polymer substrates, well-ordered SAM formation cannot effectively proceed, since the polar-functional groups formed on a polymer surface through surface modification methods using plasma or UV light are generally inhomogeneous and randomly distributed [5]. Such a heterogeneous surface cannot provide adequate support for the preparation of a wellordered SAM. As an example, Table 1 summarizes the contact angles and surface free energies of polyimide surfaces covered with NH2- and CF3-terminated SAMs, and their values are compared with those of SAM-covered silicon substrates used as standards. The formation of the SAMs is performed on photooxidized polyimide and silica-covered silicon substrates, respectively. The details of the SAM preparation method are described in elsewhere [5], but the surface energies of each sample are calculated using the following

d • Ld)1/2 + (s

P • LP)1/2 + (s

h • Lh)1/2]

substrates are focused, and their potential applications is also demonstrated.

**2. Chemical modification of a polymer surface**

L(1+cosSL) = 2[(s

**2.1 Motivation**

through solution processes.

**2.2 SAM formation**

equation:

where L is the liquid's surface energy, and SL is its contact angle on the solid surface. The superscripts d, p, and h indicate the dispersion, polar and hydrogen bonding forces, respectively. Since the L, Ld, LP, and Lh values of the three liquids are known (i.e., for water, L = 72.8 mN/m, Ld= 29.1 mN/m, LP = 1.3 mN/m, and Lh = 42.4 mN/m; for methylene iodide, L = 50.8 mN/m, Ld= 46.8 mN/m, LP = 4.0 mN/m, and Lh = 0 mN/m; and for *n*-decane, L = 23.9 mN/m, Ld= 23.9 mN/m, LP = 0 mN/m, and Lh = 0 mN/m), the surface free energy (S) of the sample is estimated to be the sum of the calculated Sd, SP, and Sh values.

Table 1. Surface free energies of SAM-covered photooxidized polyimide (PI), i.e., PIox, and are compared with those of standard, i.e., SAM-covered SiO2/Si.

Let's see the impact direct formation of FAS-SAM onto PIox, i.e., photooxidized polyimide, will have its intermolecular structure. The water contact angle of the FAS/PIox surface is somewhat smaller than that of the standard sample, that is, FAS-SAM/SiO2/Si, but high hydrophobic character can be seen even on it. However, this small difference in surface hydrophobicity influences strongly on the crystal growth behavior of metal oxide film, as described later. The reason why small difference in wetting property occurs between the surfaces can be well understood by comparison of surface free energy of the sample with that of the standard. As shown in Table 1, the S value of the FAS/PIox substrate (27.4 mN/m) is twice larger than that of a FAS/SiO2/Si substrate (11.9 mN/m). This indicates that the molecular density of the FAS-SAM attached to the PIox surface is considerably low compared to the reference substrate. In particular, the SP value of the FAS/PIox substrate (14.2 mN/m) is much larger than that of the FAS/SiO2/Si substrate (0.1 mN/m). This clearly demonstrates that the PIox surface underneath the FAS layer is probably exposed to air to some extent. It can be concluded that the degree of surface coverage by the FAS molecules differ greatly between the PIox and SiO2/Si substrates. In the sample covered with an amine group, its static-water contact angle is estimated about 69.6º which is slightly larger than that of the reference sample. This small increase can be also discussed comparing surface free energies between sample and reference. The S value (45.2 mN/m) is approximately three-fourths of that calculated for the standard, and , is slightly larger than that of a surface terminated with CH2 groups (36 mN/m). Both Sd and Sh values are consistent with those of standard, but the SP value of the AHAPS/PIox substrate (13.1 mN/m) is almost half of those obtained for the AHAPS/SiO2/Si substrates (26.2 mN/m). These experimental results suggest that the degree of surface molecular coverage is much less than those on the reference surfaces. As a result, such heterogeneous surfaces cannot provide adequate support for the preparation of well-ordered SAMs. Chemical templates consisting of such disordered SAMs are considered to be unfavorable for the deposition of inorganic materials on polymer substrates. Furthermore, after a relatively short time, such disordered moieties will lose their functionalities, resulting in a decrease of surface activity on the modified polymeric surface. Thus, to avoid hydrophobic recovery, SAM fabrication must be conducted immediately after surface photooxidation. Otherwise it is difficult to form covalent bonds between the SAM and the polymer surface, resulting in weak adhesion. Thus, in spite of polymer substrates offering the attractive advantages described above, such obstacles may prevent the preparation of high quality SAMs.

To overcome this shortcoming, this section demonstrates the formation of a well-ordered SAMs on PI sheets through the chemisorption of organosilane molecules at the solid/vapor interface. PI has been utilized primarily as a substrate for printed circuit boards due to its excellent thermal stability, mechanical strength and insulating properties. In the present study, in order to give silica-like properties to the PI surface, we fabricated an ultrathin buffer layer on the substrates using a either one of precursor vapors of tetraethoxysilane (TEOS) or 1,3,5,7-tetramethylcyclotetrasiloxane (TMCTS) prior to forming the SAM. Let's focus on the effects of this buffer layer on the quality of the obtained SAMs. Well-ordered SAM structures are formed on PI surfaces according to the process illustrated in Figure 2. In the first step Figure 8a), a commercial PI sheet is photochemically hydrophilized to give a PIox surface. Due to this photoirradiation, the PI surface becomes completely hydrophilic with its water-contact angle dropping from about 71° to 5° or less. In the next step (Figure 2b), either one of organosilane molecules is chemisorbed onto the PIox surface. The PIox sheet is placed together with a glass cup filled with 0.2 cm3 TMCTS (or TEOS) into a 65 cm3 TeflonTM container in a dry N2 atmosphere with less than 10% relative humidity. The container is sealed with a cap and heated in an oven maintained at 80°C for 3 h. Next, in order to convert the chemisorbed TMCTS (or TEOS) molecule into an ultrathin SiO2 layer, the sample is again irradiated for 30 min in a manner similar to the PI surface modification mentioned above. Finally, a SAM is formed on the SiO2-covered PIox surface through a vapor phase (Figure 2e). The thickness of the SiO2 nanolayer formed on the PIox substrate is estimated from a cross-sectional image acquired by HR-TEM, as shown in Figure 3. As clearly indicated by the arrows, an extremely thin layer is found to form on the PIox substrate, as a black contrasted image. Two important things can be seen in this crosssectional view. One is that the ultrathin layer uniformly deposits over the entire surface of the PIox. There are no defects such as a void and a pinhole. The other is an actual thickness of this layer which is estimated from this image to be only 1 nm or less. As confirmed by XPS, the binding energy of the Si2p spectrum for this sample is located at 103.6 eV, the value of which is consistent with that of a Si substrate covered with native oxide (SiO2/Si, 103.6 eV). Therefore, the nanolayer undoubtedly formed on the PIox substrate. Due to this SiO2 nanolayer coating, the surface chemical properties of the PI substrate are expected to be the same as those of SiO2/Si substrates, as evidenced in Table 2. In the case of the

surface free energies between sample and reference. The S value (45.2 mN/m) is approximately three-fourths of that calculated for the standard, and , is slightly larger than that of a surface terminated with CH2 groups (36 mN/m). Both Sd and Sh values are consistent with those of standard, but the SP value of the AHAPS/PIox substrate (13.1 mN/m) is almost half of those obtained for the AHAPS/SiO2/Si substrates (26.2 mN/m). These experimental results suggest that the degree of surface molecular coverage is much less than those on the reference surfaces. As a result, such heterogeneous surfaces cannot provide adequate support for the preparation of well-ordered SAMs. Chemical templates consisting of such disordered SAMs are considered to be unfavorable for the deposition of inorganic materials on polymer substrates. Furthermore, after a relatively short time, such disordered moieties will lose their functionalities, resulting in a decrease of surface activity on the modified polymeric surface. Thus, to avoid hydrophobic recovery, SAM fabrication must be conducted immediately after surface photooxidation. Otherwise it is difficult to form covalent bonds between the SAM and the polymer surface, resulting in weak adhesion. Thus, in spite of polymer substrates offering the attractive advantages described above, such

To overcome this shortcoming, this section demonstrates the formation of a well-ordered SAMs on PI sheets through the chemisorption of organosilane molecules at the solid/vapor interface. PI has been utilized primarily as a substrate for printed circuit boards due to its excellent thermal stability, mechanical strength and insulating properties. In the present study, in order to give silica-like properties to the PI surface, we fabricated an ultrathin buffer layer on the substrates using a either one of precursor vapors of tetraethoxysilane (TEOS) or 1,3,5,7-tetramethylcyclotetrasiloxane (TMCTS) prior to forming the SAM. Let's focus on the effects of this buffer layer on the quality of the obtained SAMs. Well-ordered SAM structures are formed on PI surfaces according to the process illustrated in Figure 2. In the first step Figure 8a), a commercial PI sheet is photochemically hydrophilized to give a PIox surface. Due to this photoirradiation, the PI surface becomes completely hydrophilic with its water-contact angle dropping from about 71° to 5° or less. In the next step (Figure 2b), either one of organosilane molecules is chemisorbed onto the PIox surface. The PIox sheet is placed together with a glass cup filled with 0.2 cm3 TMCTS (or TEOS) into a 65 cm3 TeflonTM container in a dry N2 atmosphere with less than 10% relative humidity. The container is sealed with a cap and heated in an oven maintained at 80°C for 3 h. Next, in order to convert the chemisorbed TMCTS (or TEOS) molecule into an ultrathin SiO2 layer, the sample is again irradiated for 30 min in a manner similar to the PI surface modification mentioned above. Finally, a SAM is formed on the SiO2-covered PIox surface through a vapor phase (Figure 2e). The thickness of the SiO2 nanolayer formed on the PIox substrate is estimated from a cross-sectional image acquired by HR-TEM, as shown in Figure 3. As clearly indicated by the arrows, an extremely thin layer is found to form on the PIox substrate, as a black contrasted image. Two important things can be seen in this crosssectional view. One is that the ultrathin layer uniformly deposits over the entire surface of the PIox. There are no defects such as a void and a pinhole. The other is an actual thickness of this layer which is estimated from this image to be only 1 nm or less. As confirmed by XPS, the binding energy of the Si2p spectrum for this sample is located at 103.6 eV, the value of which is consistent with that of a Si substrate covered with native oxide (SiO2/Si, 103.6 eV). Therefore, the nanolayer undoubtedly formed on the PIox substrate. Due to this SiO2 nanolayer coating, the surface chemical properties of the PI substrate are expected to be the same as those of SiO2/Si substrates, as evidenced in Table 2. In the case of the

obstacles may prevent the preparation of high quality SAMs.

SAM/nanolayer/PIox sample, a highly oriented SAM is expected to have grown similarly on the substrate. Unlike SAM sample without SiO2 nanolayer (see Table 1), the S values of the SAM-covered PIox samples (12.4 mN/m for FAS and 62.8 mN/m for AHAPS) are almost consistent with those of the standard substrate (11.9 mN/m for FAS and 61.9 mN/m for AHAPS). Specifically, the SAM formation on the nanolayer leads to the SP which nearly corresponds to the value on the reference sample. As a result, this nanolayer opens SAM chemistry even on polymer sheet [5,6].

Table 2. Surface free energies of SAMs on SiO2 nanolayer/PIox substrates, and are compared with those of standard, i.e., SAM-covered SiO2/Si.

Fig. 2. Schematic diagram for SAM formation on PIox substrate with a 1 nm thick SiO2 buffer layer.

Fig. 3. A cross-sectional HR-TEM image of TEOS-covered PIox substrate after photooxidation. The TEOS layer is converted into 1 nm thick silica layer as a consequence of the photooxidation with vacuum UV light at 103 Pa.

#### **2.3 Attachment of ceramics film on polyimide**

This subsection describes the potentials as chemical templates of SAMs formed on nanolayers in the deposition of metal and ceramics from liquid phase. To achieve a better understanding of the potentials, SAM-covered samples with and without silica nanolayer are micropatterned by photoirradiation with UV lamp through photomasks.

Among the various types of reactive functional groups, amino groups are of particular interest owing to their chemical reactivities. Palladium colloid is particularly important as a catalyst for the fabrication of metal micropatterns on non-conductive substrates like a polymer. Figure 4 illustrates a deposition process of copper microlines on polymer sheets with and without a silica nanoalyer. Each of SAM-covered PIox sheets are photolithographically micropatterned (Figure 4a). The micropatterned sample substrates are activated for 20~30 min by immersion in a Pd(II) solution containing PdCl2 and HCl (adjusted to pH 5.0 by adding aq-NaOH solution), followed by gentle rinsing with Milli-Q water (Figure 4c). These activated substrates are then immediately immersed in a commercial plating bath containing formaldehyde as a reducing agent (Figure 4d). The pH value of this solution is 12.5. As a result, copper microlines are deposited on the AHAPS chemical templates, as shown in Figure 4e. The comparison of LFM images (a) with (b) in Figure 5 shows the significant difference in a quality of SAMs formed on PIox with (b) and without (a) nanolayer. In these LFM images, bright and dark areas correspond to photoirradiated and unirradiated regions, respectively. As can be seen in Figure 5b, the 10 m-wide micropatterns are clearly imaged through the friction force contrast between the photoirradiated and masked regions of the AHAPS/nanolayer/PIox substrate. The photoirradiated regions exhibit stronger friction forces than the unirradiated AHAPScovered regions. Due to UV irradiation, an AHAPS-SAM is removed through oxidation with UV-activated oxygen species. Accordingly, in the UV-irradiated regions of the sample, the bottom part of the AHAPS layer, which consisted of a 0.2~0.27 nm-thick SiO2 monolayer,

Fig. 3. A cross-sectional HR-TEM image of TEOS-covered PIox substrate after

are micropatterned by photoirradiation with UV lamp through photomasks.

the photooxidation with vacuum UV light at 103 Pa.

**2.3 Attachment of ceramics film on polyimide**

photooxidation. The TEOS layer is converted into 1 nm thick silica layer as a consequence of

This subsection describes the potentials as chemical templates of SAMs formed on nanolayers in the deposition of metal and ceramics from liquid phase. To achieve a better understanding of the potentials, SAM-covered samples with and without silica nanolayer

Among the various types of reactive functional groups, amino groups are of particular interest owing to their chemical reactivities. Palladium colloid is particularly important as a catalyst for the fabrication of metal micropatterns on non-conductive substrates like a polymer. Figure 4 illustrates a deposition process of copper microlines on polymer sheets with and without a silica nanoalyer. Each of SAM-covered PIox sheets are photolithographically micropatterned (Figure 4a). The micropatterned sample substrates are activated for 20~30 min by immersion in a Pd(II) solution containing PdCl2 and HCl (adjusted to pH 5.0 by adding aq-NaOH solution), followed by gentle rinsing with Milli-Q water (Figure 4c). These activated substrates are then immediately immersed in a commercial plating bath containing formaldehyde as a reducing agent (Figure 4d). The pH value of this solution is 12.5. As a result, copper microlines are deposited on the AHAPS chemical templates, as shown in Figure 4e. The comparison of LFM images (a) with (b) in Figure 5 shows the significant difference in a quality of SAMs formed on PIox with (b) and without (a) nanolayer. In these LFM images, bright and dark areas correspond to photoirradiated and unirradiated regions, respectively. As can be seen in Figure 5b, the 10 m-wide micropatterns are clearly imaged through the friction force contrast between the photoirradiated and masked regions of the AHAPS/nanolayer/PIox substrate. The photoirradiated regions exhibit stronger friction forces than the unirradiated AHAPScovered regions. Due to UV irradiation, an AHAPS-SAM is removed through oxidation with UV-activated oxygen species. Accordingly, in the UV-irradiated regions of the sample, the bottom part of the AHAPS layer, which consisted of a 0.2~0.27 nm-thick SiO2 monolayer, remains on the ONS/PIox substrate surface. This SiO2 monolayer-covered nanolayer/PIox surface is highly hydrophilic with a water-contact angle of 5° or less. It therefore adheres more strongly to the Si probe surface, which is also covered with a hydrophilic native SiO2 layer. The friction force contrast on the UV-irradiated regions is thus greater than that on the AHAPS-covered areas. Although it can be expected that this SiO2 monolayer would similarly remain on the UV-irradiated regions of the AHAPS/PIox substrate, the friction force contrast in its LFM image (Figure 5a) is relatively small compared to that of the AHAPS/nanolayer/PIox substrate. In addition, several dark areas can be observed, even in the SiO2 monolayer-covered regions. This reveals that areas of weak friction exist on the

Fig. 4. Schematic illustration showing the fabrication of copper microlines on PIox substrate.

Fig. 5. LFM images of micropatterned PIox substrates covered with (a) AHAPS-SAM and (b) AHAPS-SAM/nanolayer, and AFM images and their cross-sectional views of copper microlines area-selectively electroless plated on the micropatterned (c) AHAPS-SAM/PIox and (d) AHAPS-SAM/nanolayer/PIox substrates.

SiO2 monolayer-covered regions. These are believed to be due to incomplete termination with AHAPS molecules. As mentioned in the previous section, since the packing density of the AHAPS molecules on the PIox substrate is considered to be low, parts of the PIox substrate surface which does not react with AHAPS molecule probably remain even after its SAM formation. However, the polar-functional groups terminating such a surface are unstable and gradually lose their polarity due to hydrophobic recovery, resulting in a decrease in hydrophilicity of the PIox surface. The dark regions are considered to be hydrophobic areas regenerated by this hydrophobic recovery.

Copper films are deposited on both micropatterned samples, since both the amount of Pd catalyst adsorbed and the subsequent growth of Cu through electroless plating are thought to be related to the termination states of the NH2 groups in the AHAPS molecules. Typical AFM images of the electroless-plated copper microlines and their schematic cross-sections on both micropatterned AHAPS/PIox and AHAPS/nanolayer/PIox substrates are shown in Figures 5c and d, respectively. In both images, continuous film-like deposits of Cu can be seen on the masked regions, while the surrounding photoirradiated regions remain free of deposits. 10-m-wide Cu microlines are successfully formed in a 10×10 m2 area on both substrates. The film thicknesses as estimated from these images are about 37 and 165 nm for Figures 5c and d, respectively. Accordingly, the growth rate of each is calculated to be about 3.7 and 16.5 nm/min, respectively, because of the deposition time of a 10 min in both cases. It is clear that the electroless Cu plating proceeded more efficiently on the AHAPS/nanolayer/PIox substrate. Its Cu growth rate is roughly consistent with that observed on an AHAPS/SiO2/Si substrate (20 nm/min or less). In spite of the fact that immersion time is identical, the growth rate on the AHAPS/PIox substrate without the silica buffer layer is only about one-fifth of that on the AHAPS/nanolayer/PIox substrate, offering further proof that the degree of surface coverage by Pd catalyst particles is much higher on the AHAPS/nanolayer/PIox substrate than on the AHAPS/PIox substrate. As with the plating rate, the pattern resolution also depends on the quality of the AHAPS-SAM. As clearly seen in the cross-section of the Cu microlines on the AHAPS/nanolayer/PIox substrate (Figure 5d), the edge regions are relatively sharp. On the contrary, the microlines' edge sharpness is low on the AHAPS/PIox substrate (Figure 5c). These results agree well with the LFM results described above. It is thus concluded that the silica interlayer proves effective in fabricating a well-defined AHAPS-SAM even on a polymer substrate, in a manner similar to that seen on a SiO2/Si substrate. In addition, the micropatterned substrate prepared by the present approach serves as an ideal template for spatially defined electroless Cu plating on nonconductive polymer substrate.

Next, the role of a high quality chemical template for depositing ceramics film, i.e., tin oxide, is demonstrated. The FAS-covered samples with and without silica nanolayer are photolithographycally micropatterned as shown in Figure 6a. Each patterned samples is immersed into a 1 M HCl solution containing 0.05 mol/l of SnCl2·2H2O (see Figure 6b). After immersion, the substrates are immediately washed with Milli-Q water several times, blown dry with N2 gas, and subsequently sonicated for 0.5~2.0 h in absolute toluene (Figure 6c) to remove the deposits from FAS-covered regions. We named this "solution lithography" [1]. As a result, a microstructured tin oxide film is obtained (Figure 6d). Typical optical micrographs of the films deposited on both the micropatterned FAS/nanolayer/PIox and FAS/PIox substrates (deposition time of 24 h; sonication time of 2 h) are shown in Figures 7a and 7b, respectively. As can be seen in Figure 7a, tin oxide deposits with continuous film-like features remain on the photoirradiated regions, while the

SiO2 monolayer-covered regions. These are believed to be due to incomplete termination with AHAPS molecules. As mentioned in the previous section, since the packing density of the AHAPS molecules on the PIox substrate is considered to be low, parts of the PIox substrate surface which does not react with AHAPS molecule probably remain even after its SAM formation. However, the polar-functional groups terminating such a surface are unstable and gradually lose their polarity due to hydrophobic recovery, resulting in a decrease in hydrophilicity of the PIox surface. The dark regions are considered to be

Copper films are deposited on both micropatterned samples, since both the amount of Pd catalyst adsorbed and the subsequent growth of Cu through electroless plating are thought to be related to the termination states of the NH2 groups in the AHAPS molecules. Typical AFM images of the electroless-plated copper microlines and their schematic cross-sections on both micropatterned AHAPS/PIox and AHAPS/nanolayer/PIox substrates are shown in Figures 5c and d, respectively. In both images, continuous film-like deposits of Cu can be seen on the masked regions, while the surrounding photoirradiated regions remain free of deposits. 10-m-wide Cu microlines are successfully formed in a 10×10 m2 area on both substrates. The film thicknesses as estimated from these images are about 37 and 165 nm for Figures 5c and d, respectively. Accordingly, the growth rate of each is calculated to be about 3.7 and 16.5 nm/min, respectively, because of the deposition time of a 10 min in both cases. It is clear that the electroless Cu plating proceeded more efficiently on the AHAPS/nanolayer/PIox substrate. Its Cu growth rate is roughly consistent with that observed on an AHAPS/SiO2/Si substrate (20 nm/min or less). In spite of the fact that immersion time is identical, the growth rate on the AHAPS/PIox substrate without the silica buffer layer is only about one-fifth of that on the AHAPS/nanolayer/PIox substrate, offering further proof that the degree of surface coverage by Pd catalyst particles is much higher on the AHAPS/nanolayer/PIox substrate than on the AHAPS/PIox substrate. As with the plating rate, the pattern resolution also depends on the quality of the AHAPS-SAM. As clearly seen in the cross-section of the Cu microlines on the AHAPS/nanolayer/PIox substrate (Figure 5d), the edge regions are relatively sharp. On the contrary, the microlines' edge sharpness is low on the AHAPS/PIox substrate (Figure 5c). These results agree well with the LFM results described above. It is thus concluded that the silica interlayer proves effective in fabricating a well-defined AHAPS-SAM even on a polymer substrate, in a manner similar to that seen on a SiO2/Si substrate. In addition, the micropatterned substrate prepared by the present approach serves as an ideal template for spatially defined

Next, the role of a high quality chemical template for depositing ceramics film, i.e., tin oxide, is demonstrated. The FAS-covered samples with and without silica nanolayer are photolithographycally micropatterned as shown in Figure 6a. Each patterned samples is immersed into a 1 M HCl solution containing 0.05 mol/l of SnCl2·2H2O (see Figure 6b). After immersion, the substrates are immediately washed with Milli-Q water several times, blown dry with N2 gas, and subsequently sonicated for 0.5~2.0 h in absolute toluene (Figure 6c) to remove the deposits from FAS-covered regions. We named this "solution lithography" [1]. As a result, a microstructured tin oxide film is obtained (Figure 6d). Typical optical micrographs of the films deposited on both the micropatterned FAS/nanolayer/PIox and FAS/PIox substrates (deposition time of 24 h; sonication time of 2 h) are shown in Figures 7a and 7b, respectively. As can be seen in Figure 7a, tin oxide deposits with continuous film-like features remain on the photoirradiated regions, while the

hydrophobic areas regenerated by this hydrophobic recovery.

electroless Cu plating on nonconductive polymer substrate.

Fig. 6. Schematic illustration showing the microfabrication process of metal oxide film on PIox substrate.

Fig. 7. Optical (a and b) and DFM images (c and d) of tin oxide film microstructured on two different substrates: (a), (c) on FAS/nanolayer/PIox substrate, while (b), (d) on FAS/PIox substrate, directly.

surrounding regions, which are not photoirradiated and are thus FAS-covered, are free of deposits. This clearly indicates that the deposits grown on the UV-irradiated regions stay tightly fix to the surface, while the deposits on the FAS-covered regions have been completely eliminated. On the other hand, no microstructured features can be seen on the micropatterned FAS/PIox substrate shown in Figure 7b. Both sample surfaces are observed in detail by DFM. As clearly seen in the DFM image in Figure 7c, highly resolved micropatterns composed of tin oxide film grow site-selectively on the photoirradiated regions of the micropatterned FAS/nanolayer/PIox substrate. In contrast, it is difficult to achieve clear micropatterns on the FAS/PIox substrate, as shown by Figure 7d. This sample has been treated in absolute toluene for 30 min. As can be clearly seen, the tin oxide remains deposited relatively evenly on the FAS-covered regions and there is marked distortion of the deposits on the UV-irradiated regions. When sonication is prolonged up to 2 h (image not shown), the micropattern becomes hardly recognizable by DFM. This significant difference in surface morphology between the samples must be considered in terms of the surface chemical properties of the UV-irradiated and the unirradiated regions on both samples. However, since it is difficult to determine the actual surface properties of these regions selectivity, both FAS/nanolayer/PIox and FAS/PIox substrates before and after VUV irradiation without micropatterning are characterized.

Figure 8 shows the variation in the -potential with the pH of UV-irradiated FAS/nanolayer/PIox, (i.e., SML (silica monolayer)/ONS/PIox) and FAS/PIox, (i.e., SML/PIox) substrates (indicated by the open and solid circles, respectively). As a control experiment, an observation of a SiO2/Si substrate is also performed in the same manner (indicated by the solid triangles). Over the pH range from 3 to 11, the SiO2/Si surface shows negative potentials from ca. –22 to –82 mV due to the partial ionization of surface silanol (Si-OH) groups. The isoelectric point (IEP) of this surface is estimated to pH 2.0. It is noteworthy that the -potential vs pH curves for the SML/ONS/PIox and SiO2/Si substrates are nearly identical in shape and magnitude over the entire pH range. Thus, it can be concluded that the UV-irradiated FAS/nanolayer/PIox surface shows chemical properties very similar to the SiO2/Si substrate, and that the oxide layers (i.e., the SML and nanolayer) has covered the PI substrate almost completely. On the other hand, the UV-irradiated FAS/PIox surface is found to 15-39% less charged than the SML/nanolayer/PIox and SiO2/Si substrates. In addition, the IEP of this sample, observed to be around pH 3.5, is higher than those of the SML/nanolayer/PIox and SiO2/Si substrates. This clearly indicates that the SML insufficiently covered the PIox surface. Since the density of the FAS molecules is considered to have been low, a very small part of the hydrophilic PIox surface is probably exposed to air. As a result, the molecules at the bottom of the FAS comprising SML are loosely packed. This marked difference in surface properties must account for both the final morphology and the adhesion properties of the deposits on the UV-irradiated regions.

Fig. 8. Comparison of -potentials as a function of pH for an irradiated FAS/nanolayer/PIox (open circles), an irradiated FAS/PIox (solid circles), and a silicon covered with native oxide (solid triangles), respectively. SML means siloxane monolayer.

shown), the micropattern becomes hardly recognizable by DFM. This significant difference in surface morphology between the samples must be considered in terms of the surface chemical properties of the UV-irradiated and the unirradiated regions on both samples. However, since it is difficult to determine the actual surface properties of these regions selectivity, both FAS/nanolayer/PIox and FAS/PIox substrates before and after VUV

Figure 8 shows the variation in the -potential with the pH of UV-irradiated FAS/nanolayer/PIox, (i.e., SML (silica monolayer)/ONS/PIox) and FAS/PIox, (i.e., SML/PIox) substrates (indicated by the open and solid circles, respectively). As a control experiment, an observation of a SiO2/Si substrate is also performed in the same manner (indicated by the solid triangles). Over the pH range from 3 to 11, the SiO2/Si surface shows negative potentials from ca. –22 to –82 mV due to the partial ionization of surface silanol (Si-OH) groups. The isoelectric point (IEP) of this surface is estimated to pH 2.0. It is noteworthy that the -potential vs pH curves for the SML/ONS/PIox and SiO2/Si substrates are nearly identical in shape and magnitude over the entire pH range. Thus, it can be concluded that the UV-irradiated FAS/nanolayer/PIox surface shows chemical properties very similar to the SiO2/Si substrate, and that the oxide layers (i.e., the SML and nanolayer) has covered the PI substrate almost completely. On the other hand, the UV-irradiated FAS/PIox surface is found to 15-39% less charged than the SML/nanolayer/PIox and SiO2/Si substrates. In addition, the IEP of this sample, observed to be around pH 3.5, is higher than those of the SML/nanolayer/PIox and SiO2/Si substrates. This clearly indicates that the SML insufficiently covered the PIox surface. Since the density of the FAS molecules is considered to have been low, a very small part of the hydrophilic PIox surface is probably exposed to air. As a result, the molecules at the bottom of the FAS comprising SML are loosely packed. This marked difference in surface properties must account for both the final morphology and the

Fig. 8. Comparison of -potentials as a function of pH for an irradiated FAS/nanolayer/PIox (open circles), an irradiated FAS/PIox (solid circles), and a silicon covered with native oxide

irradiation without micropatterning are characterized.

adhesion properties of the deposits on the UV-irradiated regions.

(solid triangles), respectively. SML means siloxane monolayer.

SAM formation technique demonstrated here offers the advantage that, by selecting an organosilane precursor, a well-ordered SAM with a wide variety of terminal-end groups can be fabricated on a mechanically-flexible polymer sheet (not only PI) without any marked change in surface morphology or bulk properties. Thus I expect this technique to be readily and widely applicable to novel polymer surface modification based on organosilane SAM chemistry. In addition, patterned SAM-covered polymer substrates could be applied as microtemplates for the site-selective immobilization of various kinds of materials including biomolecules, semiconductors and other functional molecules to fabricate novel polymerbased flexible devices.

## **3. Parallel detection of biomolecular recognition event on a microarray of monolayers**

Direct-writing using molecular ink plays a leading role in the development of bottom-up nanofabrication, due to its unparalleled, lithography-free properties, that allow precise positioning of the ink on a predefined surface site to obtain a high-resolution pattern on a substrate. To date, a number of direct-writing methods including microcontact printing, dippen nanolithography (DPN), microfluidic devices, nanoscale pipetting, and inkjet printing have been developed. A combination with either one of the molecular inks including SAMs, oligonucleotides, proteins, carbohydrates, nanoparticles, and polymers leads to the fabrication of its fine pattern. Many works in SAM chemistry have used the high resolution SAM patterns in particular, as chemical templates that serve as prospective candidates to miniaturize and simplify traditional laboratory techniques, bioassays, cell culture experiments, and multifunctional gas-sensing array. Achieving these goals requires the production of a chemical template in which different varieties of SAMs are precisely positioned on predefined surface sites on a given substrate. Thus far, a SAM system using a combination of thiol-ink and gold-substrate, i.e., ink-paper chemistry, has already satisfied this requirement. Actually, Mirkin and co-workers produced a microarray of four different alkanethiol SAMs by DPN [7]. The chemical functionalization of the gold surface allows the use of surface plasmon resonance and an oxidation-reduction potentiometer to detect the chemical or biomolecular interaction on thiol-template. However, its high-conductivity and optical inflexibility in the visible region has often limited the enviable versatility of SAM microarrays. This section thus describes the novel methodology to produce SAM microarrays on other substrates including glass, oxidized silicon, OH-terminated diamond, metal oxides, and polymers. The transparency in the visible region of glass is particularly essential to study biomolecular interactions under the microscope. Moreover, using a substrate diamond having both transparency and conductivity facilitates the electrochemical detection of specific biomolecular events.

The lithography-free method for oxide-based microarrays starts with the transfer of an alkoxysilane ink to the predefined surface position on an oxidized substrate under ambient conditions, as shown in Figure 9. A wide variety of pens including syringe and AFM tips can be lined up as candidates for painting with molecular inks. To verify the methodological potential, three different types of trialkoxysilanes, that is, triethoxysilylundecanal (TESUD), N-(6-aminohexyl)-3-aminopropyltrimethoxysilane (AHAPS), and octadecyltrimethoxysilane (OTS), are selected as model molecular inks. As shown in Figure 9, these organosilane inks are respectively transferred from the tips of syringes onto the predefined surface positions. At each deposit, silanization occurs during hydrolysis of the alkoxy groups by physisorbed

Fig. 9. Schematic illustration showing ink-paper chemistry on oxide substrate to fabricate a siloxane-based SAM microarray.

water on the hydrated surface, followed by a condensation reaction to give an interfacial covalent linkage between the silane molecule atom and the substrate. All of the analyses shown in Table 3 suggest that the alkoxysilane molecules form their one-molecular thick structures by the SAM formation process. The measurement of the water-contact angle is still a facile and useful method for estimating SAM formation. A high value of the watercontact angle for the OTS-treated sample is comparable to that of OTS-SAM prepared by chemical vapor deposition. A correspondence in the surface wetting property is observed between the AHAPS-treated sample and amine-terminated SAMs. The prepared OTS- and AHAPS-treated samples are then analyzed using atomic force microscopy (AFM) and ellipsometry. Although the ellipsometrical thickness of the AHAPS-treated sample is fairly in good agreement with that of AHAPS-SAM structure, the OTS-treated film is slightly thinner than its ideal thickness, that is, 25–28 Å. As expected in the case of small watercontact angle hysteresis ∆, which is defined as ∆ = a − r, both the samples' surface roughnesses measured by AFM at angstrom level are small, resulting in the successful formation of SAMs with or without less defective surfaces. On the other hand, the watercontact angle of TESUD-treated samples is substantially lower than that of the nonorientated TESUD-SAMs (e.g., a/r = 78.1º/76.2º). In order to elucidate the lower water contact-angles for this TESUD-treated sample, attenuated total reflection Fourier transform infrared (ATR-FTIR) spectroscopy, which is a convenient and well-established method of examining a SAM in surface chemical composition and intermolecular packing structure, is used. The frequencies of antisymmetric and symmetric methylene stretching modes are important indicators for exploring in-plane structures, including molecular ordering of alkane chains. The appearance of the a(CH2) at 2925−2930 cm−1 and the s(CH2) at 2855−2858 cm−1 suggests that the hydrocarbon chains in a SAM structure are disordered,

Fig. 9. Schematic illustration showing ink-paper chemistry on oxide substrate to fabricate a

water on the hydrated surface, followed by a condensation reaction to give an interfacial covalent linkage between the silane molecule atom and the substrate. All of the analyses shown in Table 3 suggest that the alkoxysilane molecules form their one-molecular thick structures by the SAM formation process. The measurement of the water-contact angle is still a facile and useful method for estimating SAM formation. A high value of the watercontact angle for the OTS-treated sample is comparable to that of OTS-SAM prepared by chemical vapor deposition. A correspondence in the surface wetting property is observed between the AHAPS-treated sample and amine-terminated SAMs. The prepared OTS- and AHAPS-treated samples are then analyzed using atomic force microscopy (AFM) and ellipsometry. Although the ellipsometrical thickness of the AHAPS-treated sample is fairly in good agreement with that of AHAPS-SAM structure, the OTS-treated film is slightly thinner than its ideal thickness, that is, 25–28 Å. As expected in the case of small watercontact angle hysteresis ∆, which is defined as ∆ = a − r, both the samples' surface roughnesses measured by AFM at angstrom level are small, resulting in the successful formation of SAMs with or without less defective surfaces. On the other hand, the watercontact angle of TESUD-treated samples is substantially lower than that of the nonorientated TESUD-SAMs (e.g., a/r = 78.1º/76.2º). In order to elucidate the lower water contact-angles for this TESUD-treated sample, attenuated total reflection Fourier transform infrared (ATR-FTIR) spectroscopy, which is a convenient and well-established method of examining a SAM in surface chemical composition and intermolecular packing structure, is used. The frequencies of antisymmetric and symmetric methylene stretching modes are important indicators for exploring in-plane structures, including molecular ordering of alkane chains. The appearance of the a(CH2) at 2925−2930 cm−1 and the s(CH2) at 2855−2858 cm−1 suggests that the hydrocarbon chains in a SAM structure are disordered,

siloxane-based SAM microarray.

resulting in a liquid alkane or an amorphous film. In contrast, shifts from their vibrational peak positions to the frequencies of a(CH2) and s(CH2) centered at 2917 cm−1 and 2950 cm−1 indicate the formation of a highly crystalline ordered SAM structure due to the strong hydrophobic interactions between adjacent alkyl chains. On the basis of this common monomolecular packing theory, the methylene stretching vibrations observed at 2916 cm−1 for a(CH2) and at 2849 cm−1 for s(CH2) in the sample, indicate that the internal hydrocarbon chains are completely ordered to produce the solid TESUD-SAM structure. Due to this well-ordering, the TESUD-SAM has a higher hydrophilic surface property, and is slightly thicker than the disordered structures (*d*ellips = 1.2–1.4 nm). The frequencies of a(CH2) and s(CH2) for the OTS-SAM are respectively observed at 2922 cm−1 and 2852 cm−1, which suggests that this monolayer is relatively ordered when compared with its amorphous film. In Table 3, the ATR-FTIR measurement indicates that the molecular structure of the AHAPS-treated sample is akin to a non-crystalline film, but its molecular thickness and surface wetting properties correspond approximately to the reported values (*d*ellips = 1.1−1.3 nm, a/r = 62.0±3º/58.0±2º). Due to the limited early reports about AHAPS-SAM, further comparison is restricted, but the measurements shown in Table 1 yield conclusive evidences about the formation of a single-molecule thick structure of AHAPS.

Table 3. Surface characterization of OTS-, TESUD-, and AHAPS-treated films.

Fig. 10. KPFM and cross-sectional images of the microstructured OTS-, TESU-, and AHAPS-SAMs formed on SiO2/Si substrate

As shown in Figure 10, SAMs' structures shown in Table 3 are microfabricated, and are analyzed using Kelvin probe force microscopy (KPFM), which promotes a better understanding of the difference in chemical properties between the SAM−covered and −uncovered surfaces. As expected, the coplanar microstructures of OTS/SiO2, TESUD/SiO2, and AHAPS/SiO2 are clearly imaged through the difference in surface potential between the SAM-covered and -uncovered regions, respectively.  Besides, the surface potential of the surrounding SiO2 surfaces is consistently lower in magnitude than those of every SAM-covered region. Depth distribution profiles acquired from those KPFM images show potential difference between the SAM-modified and -unmodified surfaces of about 6 mV (for TESUD/SiO2), 18 mV (for ODS/SiO2) and 37 mV (for AHAPS/SiO2), respectively. A surface potential distribution between the SAM-covered and –uncovered regions is given by following equation [8-10]:

#### ∆VSAM-SiOH = / ASAM0

where is the actual molecular-dipole moment normal to the substrate, SAM and0 are the permittivity of the SAM and the free space, respectively, and A is the area occupied by the molecule. Accordingly, a comparison of the theoretical dipole moment between two different surfaces leads to the identification of the terminal functional-group of the SAM, in addition to its molecular packing density. In order to calculate the dipole moments of the SAMs' precursors used here, their simplified molecular models are assumed, as shown in Figure 11. In the model SAMs, the molecules lie perpendicular to the substrate, and the terminal silicon atoms in their head-groups are replaced with hydrogen atoms. Since the directional vector of a dipole moment points from the positive charge source to the negative charge source, the three types of SAM-covered surfaces shown in KPFM images must have a positive potential contrast when compared with the surrounding SiO2 surface. In Figure 11, the theoretical dipole moment of the AHAPS-model is 29.91 Debye, oriented to the substrate with a 0° of tilt angle to its alkyl chain. The amorphous-like molecular structure of AHAPS-SAM does not allow the estimation of its tilt-angle by ATR-FTIR, but the direction of its theoretical dipole moment is consistently toward the substrate, regardless of the actual molecular tilt angle. The dipole moment of the OTS-model is equal to 0.85 Debye, and is tilted at 150.0º from the surface normal. This indicates that the OTS-SAM has a positive surface potential higher than the surrounding SiO2 surface. On the other hand, the dipole moment of the TESUD-model is tilted at 90º to the surface normal. The component of its dipole moment vertical to the substrate is thus almost zero. In order to explain the surface potential in contrast to TESUD/SiO2, the most probable dipole moment is determined as follows. Prior to the calculation of the molecular dipole moment, ATR-FTIR is used to estimate the tilt angle of the TESUD-SAM by measuring its dichroic ratio, given the strong dependence between the molecular orientation and the direction of dipole moment. The dichroic ratio, which is defined as the ratio of peak-intensity of the s-polarized absorbance, As−pol, to p-polarized absorbance, Ap−pol, gives information on the molecular tilt angle of a SAM. The dichroic ratio of the asymmetric methylene stretch observed in the TESUD-SAM has a value of 1.05, which is in fairly good agreement with those of the ordered OTS-SAMs prepared by Sagiv (D = 1.03 and 1.09) and Ulman (D = 1.04) [11,12]. Using the IR dichroism data, a tilt angle, tilt, is estimated to be 27º, with respect to the surface normal. When the TESUD-SAM is tilted at 27º with respect to the surface normal, the vertical component of the dipole moment is equal to 1.05 Debye oriented towards the substrate. As a result, the positively-charged surface potential of the TESUS-SAM on the SiO2 matrix is verified. These KPFM observations indicate that the microscale features of SAM surfaces shows chemical properties very similar to OTS-, TESUD-, and AHAPS-SAMs, respectively.

SAM-covered and -uncovered regions, respectively.  Besides, the surface potential of the surrounding SiO2 surfaces is consistently lower in magnitude than those of every SAM-covered region. Depth distribution profiles acquired from those KPFM images show potential difference between the SAM-modified and -unmodified surfaces of about 6 mV (for TESUD/SiO2), 18 mV (for ODS/SiO2) and 37 mV (for AHAPS/SiO2), respectively. A surface potential distribution between the SAM-covered and –uncovered regions is given

∆VSAM-SiOH = / ASAM0 where is the actual molecular-dipole moment normal to the substrate, SAM and0 are the permittivity of the SAM and the free space, respectively, and A is the area occupied by the molecule. Accordingly, a comparison of the theoretical dipole moment between two different surfaces leads to the identification of the terminal functional-group of the SAM, in addition to its molecular packing density. In order to calculate the dipole moments of the SAMs' precursors used here, their simplified molecular models are assumed, as shown in Figure 11. In the model SAMs, the molecules lie perpendicular to the substrate, and the terminal silicon atoms in their head-groups are replaced with hydrogen atoms. Since the directional vector of a dipole moment points from the positive charge source to the negative charge source, the three types of SAM-covered surfaces shown in KPFM images must have a positive potential contrast when compared with the surrounding SiO2 surface. In Figure 11, the theoretical dipole moment of the AHAPS-model is 29.91 Debye, oriented to the substrate with a 0° of tilt angle to its alkyl chain. The amorphous-like molecular structure of AHAPS-SAM does not allow the estimation of its tilt-angle by ATR-FTIR, but the direction of its theoretical dipole moment is consistently toward the substrate, regardless of the actual molecular tilt angle. The dipole moment of the OTS-model is equal to 0.85 Debye, and is tilted at 150.0º from the surface normal. This indicates that the OTS-SAM has a positive surface potential higher than the surrounding SiO2 surface. On the other hand, the dipole moment of the TESUD-model is tilted at 90º to the surface normal. The component of its dipole moment vertical to the substrate is thus almost zero. In order to explain the surface potential in contrast to TESUD/SiO2, the most probable dipole moment is determined as follows. Prior to the calculation of the molecular dipole moment, ATR-FTIR is used to estimate the tilt angle of the TESUD-SAM by measuring its dichroic ratio, given the strong dependence between the molecular orientation and the direction of dipole moment. The dichroic ratio, which is defined as the ratio of peak-intensity of the s-polarized absorbance, As−pol, to p-polarized absorbance, Ap−pol, gives information on the molecular tilt angle of a SAM. The dichroic ratio of the asymmetric methylene stretch observed in the TESUD-SAM has a value of 1.05, which is in fairly good agreement with those of the ordered OTS-SAMs prepared by Sagiv (D = 1.03 and 1.09) and Ulman (D = 1.04) [11,12]. Using the IR dichroism data, a tilt angle, tilt, is estimated to be 27º, with respect to the surface normal. When the TESUD-SAM is tilted at 27º with respect to the surface normal, the vertical component of the dipole moment is equal to 1.05 Debye oriented towards the substrate. As a result, the positively-charged surface potential of the TESUS-SAM on the SiO2 matrix is verified. These KPFM observations indicate that the microscale features of SAM surfaces shows chemical

properties very similar to OTS-, TESUD-, and AHAPS-SAMs, respectively.

by following equation [8-10]:

Fig. 11. Molecular structures and dipole moments of the model molecules for OTS, TESUD, and AHAPS-SAMs.

A SAM microarray is produced by combining the non-lithographical method and two different alkoxysilane inks. In a manner similar to that shown in Figure 9, two different molecular inks, TESUD and AHAPS, are respectively positioned under ambient conditions to create a microarray of two silane inks, as shown in Figure 12a. This sample is then treated in a manner very similar to the above-mentioned SAM formation process. Figure 12b shows a KPFM image of the resultant sample. In this surface potential image, well-defined microstructures composed of round-shaped 20 m features are observed in contrast to the SiO2 matrix. The surface potential contrasts between SAM-covered and –uncovered regions are +6 mV for TESUD/SiO2 and +38 mV for AHAPS/SiO2, respectively. The magnitude of each surface potential contrast approximately corresponds to the values observed in the single microarrays of TESUD- and AHAPS-SAMs in Figure 10. As a result, the successful fabrication of the SAM microarray, in which two different types of SAMs are positioned on each predefined surface site, is confirmed. Next, the potentiality of this SAM microarray as single-molecule thick platform for protein array is examined using protein G, which is an immunoglobulin-binding protein. An amine in protein G reacts with the outermost aldehyde group of the TESUD region on the microarray to give corresponding schiff base, which is reduced to the secondary amine by sodium cyanoborohydride. Throughout this reaction, protein G is covalently immobilized onto the TESUD-SAM surface. Due to the use of protein primarily labeled with Alexa 488, its presence on the sample surface is observed under a fluorescence microscope. As shown in Figure 12c, a strong fluorescence signal is detected with area-selectively on the TESUD-covered region. This successful detection, even after strong surface-cleaning, implies that protein G is not simply physisorbed, but is covalently bound to the aldehyde-terminated SAM surface. In contrast, the dark areas in Figure 12c correspond to the regions covered with AHAPS-SAM. This indicates that the AHAPS-SAM surface strongly prevents even nonspecific adsorption of protein G from the contacting buffer when compared with that on SiO2 matrix. Accordingly, the fluorescence signal from the Alexa-labeled protein is the highest in magnitude on the TESUD-SAM region of the microarray.

Fig. 12. (a) Optical image of a microarray having droplets of AHAPS and TESUD molecules, (b) KPFM image of TESUD/AHAPS-SAMs microarray, (c) Fluorescence microscopic image of TESUD/AHAPS-SAMs microarray after treated with Alexa-labeled Protein G. Scale bar: 20 m.

As demonstrated here, the use of SAM microarray allows the parallel detection of different protein-behaviors on one platform surface under same experimental conditions. Although the two different kinds of SAMs are microstructured in this case, further increase in varieties will allows the parallel detection of many host-guest molecular interactions. Thus, such parallel observation of different variations in chemical reaction and biomolecular interaction on the same platform, where a number of different functional SAMs are integrated at predefined positions, will be able to lead to the achievement of facile and fast medical screening, or miniaturizations of traditional laboratory techniques, bioassays, and cell culture experiments.

## **4. Highly luminescent silicon nanoparticles**

### **4.1 Background**

232 Nanofabrication

A SAM microarray is produced by combining the non-lithographical method and two different alkoxysilane inks. In a manner similar to that shown in Figure 9, two different molecular inks, TESUD and AHAPS, are respectively positioned under ambient conditions to create a microarray of two silane inks, as shown in Figure 12a. This sample is then treated in a manner very similar to the above-mentioned SAM formation process. Figure 12b shows a KPFM image of the resultant sample. In this surface potential image, well-defined microstructures composed of round-shaped 20 m features are observed in contrast to the SiO2 matrix. The surface potential contrasts between SAM-covered and –uncovered regions are +6 mV for TESUD/SiO2 and +38 mV for AHAPS/SiO2, respectively. The magnitude of each surface potential contrast approximately corresponds to the values observed in the single microarrays of TESUD- and AHAPS-SAMs in Figure 10. As a result, the successful fabrication of the SAM microarray, in which two different types of SAMs are positioned on each predefined surface site, is confirmed. Next, the potentiality of this SAM microarray as single-molecule thick platform for protein array is examined using protein G, which is an immunoglobulin-binding protein. An amine in protein G reacts with the outermost aldehyde group of the TESUD region on the microarray to give corresponding schiff base, which is reduced to the secondary amine by sodium cyanoborohydride. Throughout this reaction, protein G is covalently immobilized onto the TESUD-SAM surface. Due to the use of protein primarily labeled with Alexa 488, its presence on the sample surface is observed under a fluorescence microscope. As shown in Figure 12c, a strong fluorescence signal is detected with area-selectively on the TESUD-covered region. This successful detection, even after strong surface-cleaning, implies that protein G is not simply physisorbed, but is covalently bound to the aldehyde-terminated SAM surface. In contrast, the dark areas in Figure 12c correspond to the regions covered with AHAPS-SAM. This indicates that the AHAPS-SAM surface strongly prevents even nonspecific adsorption of protein G from the contacting buffer when compared with that on SiO2 matrix. Accordingly, the fluorescence signal from the Alexa-labeled protein is the highest in magnitude on the TESUD-SAM

Fig. 12. (a) Optical image of a microarray having droplets of AHAPS and TESUD molecules, (b) KPFM image of TESUD/AHAPS-SAMs microarray, (c) Fluorescence microscopic image of TESUD/AHAPS-SAMs microarray after treated with Alexa-labeled Protein G. Scale bar:

As demonstrated here, the use of SAM microarray allows the parallel detection of different protein-behaviors on one platform surface under same experimental conditions. Although the two different kinds of SAMs are microstructured in this case, further increase in varieties

region of the microarray.

20 m.