MICAtronics: A new platform for flexible X-tronics

FlatChem 3 (2017) 26–42

Contents lists available at ScienceDirect

FlatChem j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / fl a t c

MICAtronics: A new platform for flexible X-tronics Yugandhar Bitla a, Ying-Hao Chu a,b,c,⇑ a

Department of Materials Science and Engineering, National Chiao Tung University, Hsinchu 30010, Taiwan Institute of Physics, Academia Sinica, Taipei 11529, Taiwan c Material and Chemical Research Laboratories, Industrial Technology Research Institute, Hsinchu 31040, Taiwan b

a r t i c l e

i n f o

Article history: Received 5 May 2017 Revised 7 June 2017 Accepted 7 June 2017

Keywords: Two-dimensional materials van der Waals epitaxy Mica Flexible electronics Functional oxides

a b s t r a c t In the present era of ‘‘Internet-of-Things”, the demand for flexible, light-weight, low-cost, low-power consumption, multifunctional, and environmentally friendly electronics has moved to the forefront of materials science research. Numerous compounds with unique material properties in epitaxial thin film form hold key to future technologies. van der Waals epitaxy (vdWE) involving two-dimensional layered materials can play a crucial role in the expansion of thin film epitaxy by overcoming the bottleneck of material combinations due to lattice/thermal matching conditions inherent to conventional epitaxy. Among the layered materials, mica is a well-known phyllosilicate mineral that can have a remarkable impact on flexible electronics. We confine ourselves to the validity of vdWE of functional oxides on muscovite mica throughout this treatise. These heterostructures with excellent properties are flexible and exhibit high-temperature stability. With such demonstrations, it is anticipated that MICAtronics, vdWE on mica, can reveal unusual properties and emergent phenomena in the realm of high-performance flexible device applications. Ó 2017 Elsevier B.V. All rights reserved.

Contents Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conventional vs van der Waals epitaxy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mica: structure and properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Functional oxides on mica for flexible applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Electronics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Transparent conducting oxides (ITO, AZO) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . VO2 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . MoO2 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Spintronics (Fe3O4) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Magnetostrictive CoFe2O4 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Non-volatile memory (PZT) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Multiferroics (Self-assembled BFO-CFO nanocomposite) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Outlook and future scope . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

26 27 28 31 31 31 33 34 34 35 37 38 39 39 40 40

Introduction

⇑ Corresponding author at: Department of Materials Science and Engineering, National Chiao Tung University, Hsinchu 30010, Taiwan. E-mail address: [email protected] (Y.-H. Chu). http://dx.doi.org/10.1016/j.flatc.2017.06.003 2452-2627/Ó 2017 Elsevier B.V. All rights reserved.

The primary obstacles to success of conventional epitaxy are the lattice and thermal expansion mismatches between the overlayer and substrate materials that seriously curtail material combinations. Therefore, an alternative approach capable of overcoming

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

aforementioned constraints can expand the domain of the heteroepitaxy of high-quality three-dimensional (3D) materials further. van der Waals (vdW) epitaxy, a new paradigm of heteroepitaxy involving two-dimensional (2D) layered materials with unique and promising properties has opened new avenues for fundamental scientific studies and applied device designs [1– 12]. Prompted by the advent of graphene, the study of 2D materials [13] with novel electrical, optical, thermal, and mechanical properties has enriched this new research area and emerged as ‘‘one of the hottest topics in materials research”, due to technological payoff promise. They are often referred to as van der Waals solids due to strong intra-layer but weak inter-layer interactions. Normally, 2D layered structures are obtained via mechanical exfoliation [14], chemical isolation [15,16] and vapor deposition. Mica is a well-known layered silicate compound which is used in one form or the other in our daily routine yet very little is known about its origin, composition, and properties. For many years, natural or synthetic mica has been used in industry as an insulating material, as a dielectric in capacitors, an insulator in home electrical appliances like hot plates, toasters, and irons. Ground mica is used in roofing materials and waterproof fabric coatings, as filler in many compounds and mixtures such as paint, wallpaper, joint cement, plastics, cosmetics, well-drilling products, and a variety of agricultural products. The chemical inertness, high mechanical strength, suitable thermal and biological properties of mica have also generated interest in biomedical field [17,18]. On the research front, mica has served as an unmatched substrate for surface force measurements due to its atomically smooth surface. The surface of mica provides an interesting playground for the study of 2D ion diffusion, 2D crystal nucleation and surface conductivity phenomena. Normally, freshly cleaved mica surface is hydrophobic but after long enough hydration time, it becomes hydrophilic (wetted by water) with structured water layer close to its surface [19,20]. The first structured water layer on mica is reported to be ferroelectric [21] with O–H bonds either pointing towards mica or linking adjacent water molecules and oxygen atoms facing the air interface [22]. As mica is a common surface for imaging biological macromolecules, the ordered water has strong affinity to them [23]. It is argued that air-cleaved mica is covered by one potassium carbonate formula unit while its formation is still unclear [24]. Recently, ultrathin mica sheet was used as an insulator for large area organic thin film transistor [25–27] arrays with promising implications in organic electronics and as a substrate for the growth of functional oxides [28–36]. The main reason why the 2D materials are advantageous to be used as substrates is that the overgrown film materials float on top of the dangling-bondsfree substrate surface instead of being rigidly bound to it, thereby mitigating lattice and thermal mismatch between over layer and the underlying 2D substrate. Therefore, MICAtronics, vdW epitaxy on mica, provides new arena of multifunctional flexible X-tronics. Before envisaging mica’s applicability in flexible device applications, several issues needed to be addressed, including novel mechanics, architectures, designs, new materials and fabrication methods. In this review, a strong emphasis is laid upon the properties, integration, and applications of oxide heteroepitaxy on muscovite mica via the weak van der Waals interaction [28–34]. First, a contradistinction between conventional and vdW epitaxy is provided followed by a brief introduction to 2D layered mica central to this review along with its basic properties and preparation. Second, the nanoscale engineering of unconstrained epitaxial oxides on muscovite mica with atomically sharp interfaces and zero interdiffusion of atoms in the creation of a wide range of functionalities spanning ultrathin, flexible and transparent optoelectronic, spintronic, ferroelectric and multiferroic with suitable examples are

27

outlined. Finally, future research directions and the technological implications of MICAtronics for practical devices are highlighted.

Conventional vs van der Waals epitaxy Multitudes of compounds with unique material properties in epitaxial thin film form hold key to future technologies. Although the high crystalline quality associated with epitaxial growth plays a crucial part in realizing optimal properties, the grown layer becomes rigidly attached to its substrate. This substrate clamping effect has to be minimized, if not removed, to obtain the strainfree thin films with ideal properties of the bulk. The severity of this clamping effect is mostly controlled by the nature of the interface in the heterostructure. Depending on the strength of interfacial interactions, the thin film epitaxy is classified into (1) conventional epitaxy: strong interactions creating the reactive interfaces with chemical reaction of the constituents or interdiffusion and (2) van der Waals epitaxy: weak interactions resulting in the nonreactive and abrupt interfaces. In the conventional epitaxy, the chemical bonds in the filmsubstrate contact plane determine the lattice matching conditions and the orientation of the epilayer [37]. These chemical bonds (covalent, ionic and metallic) orient across the interface as given by the symmetry of substrate and film, and accommodate differences in lattice mismatch by interfacial stress or strain eventually modifying overlayer’s cell parameters (Fig. 1a). Normally, an abrupt interface is experimentally achieved in the epitaxy of lattice matched systems. For lattice mismatched systems, high quality pseudomorphic heterointerfaces with very thin and strained layers with changes in their electronic structure and properties or poorly controlled interfaces dominated by dislocations and defects or interdiffusion are formed. As a result, only a few selected number of lattice matched 3D heterointerfaces exhibit ideal bulk-like properties. Moreover, non-saturated ‘‘dangling” bonds in the contact plane introduce misfit dislocations during the film growth and also impart substrate clamping effect undermining the heterostructures’ performance. Therefore, such limitation restricts the heteroepitaxy of technologically important materials. A new kind of epitaxy involving weak vdW interactions between thin film-substrate yielding high-quality material with reduced threading dislocation densities [38–43] was first reported by Koma. Earlier theoretical framework [44,45] involving weaker overlayer-substrate interactions than interactions within the overlayer predicted an incommensurate epitaxy with unrelated orientation relationship between overlayer-substrate to the crystalline symmetry. Weak intermolecular/interionic coulombic interactions are the basis for vdW bonds wherein there is no electron transfer or sharing between ions or molecules participating in the bond [46]. vdW bonding energies (40–70 meV) are much smaller than covalent bonding energies (200–6000 meV) [4]. In vdW epitaxy, the epilayer grows with its bulk lattice constant despite lattice mismatch as large as 60% [47,48]. In principle, no stress or strain is induced by the substrate across the contact plane. Therefore, epilayer has bulk ideal single crystal like crystalline quality and the vdW interface is abrupt on an atomic scale which is free of intermixing and chemical cross contamination. vdW epitaxy is an incommensurate epitaxy that presents incoherent interfaces (Fig. 1b). Therefore, the prospects of vdW epitaxy as an alternative route are excellent towards the successful heteroepitaxial growth of planar as well as nonplanar (nanowire arrays, tripods, tetrapods and complex nanostructures) geometries. It is instructive that vdW epitaxy can expand the domain of thin film heteroepitaxy beyond the strain engineering phase space of limited material combinations permitted by the conventional epitaxy.

28

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

Fig. 1. Schematic illustration of conventional and van der Waals epitaxy.

Initially, vdW epitaxy strictly implied the growth of (2D) layered material on another (2D) layered substrate (Fig. 1b) but later it extended to the heteroepitaxial growth of (2D) layered materials onto (3D) substrates or vice versa (Fig. 1c and d). They come under the quasi-vdW epitaxy (Fig. 1c and d) which is intermediate to conventional and vdW epitaxies. Throughout this review we use vdW epitaxy for the description of both vdW and quasi-vdW epitaxial processes, ie., epitaxy involving a 2D material. The vdW epitaxy of functional oxides on the 2D layered muscovite mica concerned herein will highlight some of these combinations. As is evident from Fig. 1, the nucleation and growth in conventional 3D/3D epitaxy is completely different to vdW epitaxy. The former involves chemisorption process while the latter is governed by physisorption process. In Quasi-vdW epitaxy, initially an interfacial layer is formed which strongly interacts with the 3D material via directional bonds, onto which further film growth proceeds homoepitaxially. In principle, the interactions that decide the nucleation and morphology of the epilayer in vdW epitaxy should depend on thermodynamic and kinetic properties of the overlayer. In a thermodynamic approximation, 2D triangular nucleation clusters result when surface and interface tensions are dominant during vdW growth [47]. Experimentally, it is a daunting task to find the right conditions for the vdW epitaxy as the weak interlayer interaction normally favours island growth rather than that of continuous monolayers [42]. Nevertheless, different material combinations characterized by a vdW-type phase boundary present different overlayer nucleation and morphologies suggesting a different degree of adatom-substrate electronic states coupling. Thus, a small but non-negligible electronic overlap across the vdW-gap (different for the different 2D materials) dependent on the orientation of the stacked sandwich units governs the bonding between the layers and thus, the azimuthal epitaxial orientation. As stated above, these bonding interactions across the vdW-surface of a layered compound that are two orders of magnitude weaker than those within the plane are weak enough to accommodate the overlayer-substrate misfit and in most cases, strong enough to induce epitaxial orientation. However, if this mechanism of vdW epitaxy were true then it must occur at all mismatched vdW heterointerfaces. Therefore, a detailed understanding of the elec-

tronic coupling across vdW-heterointerfaces is necessary to address this issue. More so, the effect of anisotropic interactions (strong intra-plane and weak inter-plane) and the fundamental mechanisms that govern nucleation and growth during vdW epitaxy are not well understood yet. Therefore, nucleation and growth processes in vdW epitaxy forms an interesting subject for further detailed investigation.

Mica: structure and properties Micas are the most abundant subgroup of phyllosilicates that exhibit sheet/layered structure. The crystal structure of mica was reported [49,50] in 1920s and the ideal unit cell formula can be represented by X2YnZ8O20(OH,F)4 where X represents interlayer cations (K+, Na+, Ca2+), Y represents the octahedral coordinated elements (Al3+, Mg2+, Fe2+, Li+,..) and Z represents the tetrahedral coordinated elements (Si4+, Al3+,..). They are further classified into dioctahedral (n = 4) such as muscovite, pargasite, etc., and trioctahedral (n = 6) such as pholgopite, biotitite, zinwaldite, etc. Thin film fabrication processes employ muscovite [KAl2(Si3Al) O10(OH)2: a = 5.17 Å, b = 8.94 Å, c = 20.01 Å, and b
96°] [50], biotite [K(Mg, Fe)3(AlSi3O10)(F, OH)2], and phlogopite [KMg3(AlSi3O10) (OH)2] as the common substrates. All are monoclinic with a tendency towards pseudo-hexagonal structure on their basal plane. The layered mica structural unit (
1 nm) comprises of two tetrahedral (T) sheets on either sides of an octahedral (O) sheet and these (2:1) layer stacks are bound together by interlayer cations (Fig. 2a). Silicate tetrahedron (SiO4) and aluminium octahedron (AlO6) are the basic building blocks of these layers. The tetrahedral sheet consists of honeycomb arrangement of hexagons of SiO4 tetrahedra, where each tetrahedron shares its basal oxygens with the adjacent tetrahedron. Si ions of these layers are partially substituted by Al ions (1/3) to give a net negative charge. The apical oxygens together with hydroxyl group (OH) lying below the center of SiO4 hexagon of tetrahedral layer form the octahedral layer. The center of each octahedron is occupied by Al (muscovite mica) or Mg (pholgopite mica). Below the octahedral layer is a second tetrahedral layer. This T-O-T stack is held together by interlayer cations (K+) maintaining the charge neutrality of the unit. Fig. 2b shows

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

29

Fig. 2. Crystal Structure. (a) Structural unit of mica. Surface structure of muscovite mica: (b) [001] projection and (c) [100] projection.

the basal plane [001] of the cleaved muscovite mica surface with a hexagonally arrayed pattern of oxygen atoms, with the in-plane lattice parameters marked. Fig 2c shows the [100] projection of muscovite mica where the TOT sandwich units are separated from each other by the van der Waals-gap. These sandwich units have a very strong covalent bonding within them while the sandwich units are weakly held together by K+ atoms along the crystallographic c-axis. Cleavage of mica along this vdW gap layer produces two large, atomically flat surfaces, with each surface occupied by equal but randomly distributed K+ atoms to preserve charge neutrality. The mica presents different stacking sequences which may deviate from each other by a different metal coordination and/or by a different stacking of the sandwich units along c. The following salient features make muscovite mica an ideal substrate for flexible device applications. i. 2D structure: Exhibit strong intra-layer but weak inter-layer interactions prerequisite for vdW epitaxy. Due to the weak vdW interaction, it is possible to peel off overlayer from mica creating an unconstrained or free-standing like overlayer thin film. It is more natural for the perovskite to be grown on mica with a perovskite {111} orientation since it is hexagonal and should fit better with the substrate. ii. Transparent: A thin mica (
100 mm) sheet exhibits extraordinary transmittance in the ultraviolet–visible-infrared range of the electromagnetic spectrum. They can serve as very attractive templates for a range of optoelectronic applications. iii. Elastic: The suspended mica nanosheets exhibit high Young’s modulus of Y
200 GPa, which present low prestrain (<0.25 N/m) and high breaking force (4–9 GPa) [51]. These mechanical properties of atomically thin mica crystals validate their applicability as dielectric substrates in highly demanding mechanical applications. This feature is extremely useful in flexitronics-paper electronics, wearable electronics, conformal electronics and so on.

iv. Flexible: Mica possess a very high yield strain and the thickness of a single cleaved mica sheet can be controlled down to few microns by splitting manually or using a scotch tape. Few microns thick mica sheet is flexible and a bending radius down to 0.03 cm is reported for 100 nm thick mica sheet [25]. v. Chemically inert: The absence of dangling bonds on surface make them incapable of forming chemical bonds. They are tolerant against chemicals. vi. Non-toxic: Beneficial in biomedical and wearable applications vii. Thermally conductive: The heat generated by flexible device units on muscovite facilitate effective dissipation of heat. viii. Electrically insulating: Prospective use as a substrate as well as a gate dielectric in flexible electronic devices [25]. ix. High thermal resistance: Its melting point (1150 K - 1300 K) is also high [52] to be compatible in modern thin film processes. x. CTE matched with Si: Coefficient of thermal expansion (CTE) of mica with that of Si. xi. Atomically flat surface: This is the key reason for using mica in TEM and AFM. Moreover, epilayers on mica are ideal systems for investigating molecular surface phenomena like adhesion, friction and colloidal interactions [53]. xii. Inexpensive and abundant: Phyllosilicates are one of the most abundant minerals found in earth’s crust. xiii. Non-magnetic: Natural mica prepared is expected to be diamagnetic [54] but sometimes very minute magnetic impurity inclusions like Fe can result in a weak anisotropic ferromagnetism. xiv. Light-weight. xv. Compatible with all deposition techniques. xvi. Biocompatibility: This can have significant impact on bioinspired applications [55]-artificial skin and muscles, prosthetic limbs, soft and humanoid robots, smart clothing and electronic textiles.

30

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

They are additionally hydrophyllic (wetted by water), resistant to weathering and corrosion attack by acids of alkalies. All these features qualify mica as a destination substrate for flexible electronics [56,57]. Other commonly used flexible substrates are, ultra-thin glass, polymers, paper/fibrous materials and metal foils. Glasses with thickness smaller than 0.3 mm are bendable, thermally stable and highly transparent but they are fragile and costly. Polymer substrates (polystyrene (PS), polyimide (PI), polymethylmethacrylate (PMMA), polyethylene terephthalate (PET), polycarbonate (PC), polyethersulfone (PES), polydimethylsiloxane (PDMS) and polyethylene naphthalate (PEN)) have very good mechanical compliance but they are not thermally stable and hence, hinder the growth of high quality films. Fibrous materials have attracted much attention for flexible electronic applications [58–61] because of their extremely low cost but they pose serious problems due to sensitivity to water, solvents and very rough surfaces. Metal foils are thermally stable, have excellent thermal and electrical conductivities, solvent resistance but they are nontransparent. From Table 1, it is imperative that mica exhibits superior characteristics as a flexible template for flexible applications compared to its counterparts. Flexitronics refers to all flexible device structures, in general and flexible electronics, in particular, which are tolerant to a range of mechanical deformation modes [62]-bending, stretching and twisting. Therefore, evaluating and understanding the device per-

formance under these mechanical constraints is of prime concern in this area. However, we confine ourselves to the bendable versions of planar structures. In a bent state of the thin film heterostructure (Fig. 3a), the substrate and thin film experience different types of strains. The physical strain [62,63] S imposed on a thin film layer of thickness tf located at the surface of a substrate of thickness tS by bending the substrate into a radius R is

 t þt 1þ2gþvg2 given by S ¼ f2R S ð1þ gÞð1þvgÞ where g = tf/tS, v = Yf/YS, Yf and YS are Young’s modulii of the thin film and the substrate, respectively. In a bent sheet, the greatest strains occur on the surfaces. A mechanically neutral plane (dotted line) exists within the sheet where neither tensile nor compressive strain is present but opposite type strains occur across this plane. In general, tf << tS, the imposable strain is directly proportional to the substrate thickness and inversely proportional to the bending radius. Therefore, at a given strain value, the minimum bending radius can be improved by choosing a thinner substrate. In other words, heterostructures grown on thinner substrates tend to tolerate smaller bending radii before failing. For a 100 nm mica sheet, the minimum radius of bending is 0.3 mm [25]. All the flexible heterostructures demonstrated in the following sections will be examined either in Flexin mode as shown in Fig. 3b (thin film under compression) or Flex-out mode shown in Fig. 3c (thin film under tension). The temperature-induced curvature in film-substrate biomorph due their differential thermal expansions is given by

Table 1 Flexible substrates for flexible device applications. PS

PC

PDMS

PEN

PES

PET

Poisson ratio Refractive index

0.34 1.6

0.37 1.58

0.5 1.4

0.3–0.4 1.5–1.75

0.4 1.66

0.35–0.45 0.34 1.66 1.7

Transmittance (%) Melting point (°C) Operating temp (°C) TEC (106 /K)

80–90 240 50–90 30–210

90 225 115–150 66–70

45 260 1

Electrical resistivity (Ocm) 1018–1020 1014–1016
1015 Thermal conductivity (W/m. K) 0.033 0.2 0.15 Youngs modulus (GPa) Tensile strength (MPa) Compression strength (MPa) Dielectric constant


3 46–60 70 2.6


2 55–75 >80 2.9


6  104 2.24 2.5

PI

PMMA

Glass

SS

0.35–0.4 1.49

0.238 1.5

0.3

muscovite phlogopite

0.25 na: 1.563 nb: 1.596 nc: 1.60 87–90 90 89 yellow 80–93 92 0 >90 270 >250 520 130–160 720 >1500 1000 155–180 223 115–170 250–320 50–90 600 1400 700 20,16 54 20–80 30–60 70–77 3–5 10 8–12 (||) 15–25 (\) 15 17 14 17 19 12 10 10 >10 10 >10 10 1012–1015 0.15 0.13–0.18 0.15–0.24 0.12 0.17–0.25 1.16 16 0.3 (||) 3 (\) 5–6 2.2 2–4 2–3 2–3 75 200 172 200–275 83 55–75 150–230 47–79 22–32 400 225–300 80 51 70–120 250 221 2.9 3.7 3.6 3.4 2.6 5.1 6.5–9.3

na: 1.55 nb: 1.59 nc: 1.60 >90 1100 800 5–13 (||) 0.001–1(\) 1010–1013 0.3 (||) 3 (\) 172 255–300 221 5–6

Fig. 3. Bending modes. (a) Strains experienced by substrate and thin film under bending. (b) Flex-in/compressive strain mode and (c) Flex-out/tensile strain mode of the thin film.

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

 R¼

 2 tS ð1  vg2 Þ þ 4vgð1 þ gÞ2 6ð1 þ mÞðaf  as ÞDT vg ð1 þ gÞ

!

where a is the coefficient of thermal expansion, DT represents the temperature change and m is the Poisson’s ratio. Functional oxides on mica for flexible applications Oxide materials are ubiquitous in modern science and technology. They display a huge variety of compositions and structures with interesting properties spanning ferromagnetic, ferroelectric, multiferroic, catalytic, photoluminescence, superconducting, and semiconductive behaviors. Intimately coupled and interwoven lattice, spin, charge and orbital degrees of freedom are responsible for such rich spectrum of properties. Reduced dimensionality, as an additional degree of freedom, enabled the growth of epitaxial films, multilayers, superlattices and well-ordered nanostructures (vertical and horizontal) on the nanoscale with unconventional properties relevant to various technological applications [64–66]. MICAtronics, vdW epitaxy on mica, is likely to have a significant impact on the high-performance flexible optoelectronics, sensors, memories, energy harvesting devices and so on as depicted in Fig. 4. Although, we restrict ourselves to oxide epitaxy on muscovite mica in this review, MICAtronics spectrum includes a wide variety of non-oxides too. For example, mica’s epitaxy with parahexaphenyl [67], graphene [68,69], Au [70], Ag [71], Ni [72], ZnTe [48], transition metal dichalogenide [73,74], topological insulator [75,76], P-N junction [77], solar cell [78], and so on has been realized. Nonplanar nanostructures like ZnO [79], WO3 [80] nanowires, tripods and tetrapods [48,81,82] have also been realized on mica. A vdW crystal with alternate stacking of graphene and mica is superficially speculated to result in high TC superconductivity with graphene playing the role of conducting CuO planes and the high-k dielectric mica providing the interplanar spacing [1]. To underscore the versatility of mica for flexible applications, functional oxide heteroepitaxy via pulsed laser deposition technique has been demonstrated in the following sections.

31

Electronics Transparent conducting oxides (ITO, AZO) Transparent conducting oxides (TCO) are the fundamental components in optoelectronic applications [83,84] ranging from solar cells, flat panel displays, photocatalysis, photodiodes, energy efficient windows, thin film transistors, and gas sensors. Last decade has seen dramatic technological advances in flexible, transparent and portable electronics which necessitated the development of flexible TCO. However, these demands are severely hampered by the amorphous nature of most of the flexible substrates and/or their inability to high-temperature growth causing poor device performance. Representative TCOs-Indium Tin Oxide (ITO), the best TCO so far, and Al doped ZnO (AZO), an indium-free TCO, were taken up to test the validity of vdW epitaxy on mica for the realization of future high-performance flexitronics. The high-resolution cross-sectional transmission electron microscopy (TEM) images in Fig. 5a and b taken along (a) [100]Mica and (b) [010]Mica zone axes reveal clear, defect-free and semicoherent (a) ITO/YSZ, YSZ/mica, and (b) AZO/Mica interfaces. The corresponding Fast Fourier Transform (FFT) patterns along with indexed reciprocal lattices of ITO (red), YSZ (orange), AZO (blue) and muscovite (yellow) marked in Fig. 5a and b are also shown.  The epitaxial relationships- (222)ITO||(001)Mica and [110] ITO|| [010]Mica for the ITO/mica heterostructure, and (001)AZO||(001)Mica and [010]AZO||[010]Mica for the AZO/mica heterostructure in agreement with the x-ray diffraction (XRD) results are obtained. YSZ buffer layer was chosen to realize epitaxial ITO on mica due to the excellent lattice matching (<1.6%) and cube-on-cube relationship [85,86] between YSZ and ITO. When no YSZ is used, both ITO (111) and (001) peaks were observed. The photograph of 200 nm ITO/mica, AZO/mica along with the reference mica in the inset of Fig. 5c and d clearly highlights the high optical transparency of TCO/mica heterostructures. The thickness dependent electrical conductivity of TCO/mica (Filled circles), epitaxial TCO [85–91] (Asterisk symbol), polycrystalline TCO [87,92–107] (Open circles) and TCO on flexible substrates[35,92–94,99,100,108–121] (Open circles with cross marks) compared in Fig. 5c shows that the TCO/mica has conductivity values comparable to epitaxial TCO obtained on rigid substrates. The Hall measurements confirmed the n-type semiconducting nature of the TCO films. The extracted mobility of TCO/mica as a function carrier concentration displayed in Fig. 5d are on par with epitaxial TCO as they overcome the grain boundary scattering inherent to those on flexible substrates. It is clear that the electrical properties of TCO films improve with the thickness and are comparable to the reported ones; the best parameters are achieved for
150(250) nm ITO (AZO) thick films with resistivity q
5(30) 105 Ocm and mobility l
50(20) cm2/Vs. An average visible transmittance of 68–97% (76–90%) for ITO/mica (AZO/mica) heterostructure with optical band gap of 3.7 eV (4.1 eV) was observed. The thin films less than 200 nm thickness had a very wide transparency window beyond visible region which narrowed upon increasing thickness with the effect that the transmittance and its range decreased at a faster rate in ITO/mica than in AZO/mica. The optical transmittance of TCO/mica as a function of sheet resistance presented in Fig. 5e displays higher transmittance and lower sheet resistance comparable to epitaxial TCO. Further, the figure of merit (UTC) [122],

UTC ðX1 Þ ¼ T 10 =RS , where T is the average optical transmittance

Fig. 4. MICAtronics, van der Waals epitaxy on mica, and its technological implications.

and Rs the sheet resistance, as a function of thickness is shown in Fig. 5f. It is known that the higher the UTC the better the quality of the transparent conducting films [123]. Higher figure of merit values for epi TCO and TCO/mica can attributed to their improved structural quality.

32

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

Fig. 5. TCO structure and properties. TEM image of (a) ITO/YSZ/mica and (b) AZO/mica. (c) Thickness dependent conductivity, (d) Hall mobility as a function of carrier concentration, (e) average transmittance and (f) figure of merit of TCO/mica as against epitaxial, polycrystalline and flexible counterparts. Inset shows the transparent ITO/ mica, AZO/mica along with mica substrate. Reprinted with permission from Ref. [28]. Copyright 2016 American Chemical Society.

In view of flexible optoelectronic applications, the mechanical stability of TCO/mica of different thicknesses was tested under flex-in mode and flex-out mode as depicted in Fig. 6a and b. The constant sheet resistance under bending radii down to 4 mm (7 mm) and 6 mm (9 mm) during both the compressive and tensile deformations, respectively, for ITO(AZO) are better than those reported for ITO on mica [35,75], PEN [113,124], and PET [125– 127] (AZO on PES [108]). Beyond this critical bending radius, the sheet resistance increases drastically and does not return back to

the initial value upon release to the unbent state. It is also seen that the critical bending radius of curvature decreases with TCO thickness. It is noteworthy that the critical compressive strain of TCO (ecomp = 0.75% for ITO and 0.80% for AZO) is higher than the critical C tensile strain (eten C = 0.41% for ITO and 0.46% for AZO) and are com50 33 parable to eten and ecomp = 1.7% for C = 1.1% for ITO , 1.03% for AZO C ITO [128], 1.74% for AZO [108]. Additionally, the conduction fatigue (RS as function of bending cycle) of TCO/mica in flex-in mode displayed in Fig. 6c shows that AZO/mica exhibits very stable sheet

Fig. 6. Flexible TCO/mica. Sheet resistance as a function of bending radius in (a) flex-in and (b) flex-out mode. Sheet resistance under flex-in mode as a function of (c) bending cycles and (d) time. (e) Thermal and (f) chemical stabilities. Reprinted with permission from Ref. [28]. Copyright 2016 American Chemical Society.

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

resistance even after 2000 bending cycles under a radius of curvature of 8 mm while ITO/mica exhibits a 10% increase in its resistance beyond 100 cycles. The inherent fragile nature of ITO may impose this limitation. Note that the stable conduction retention (RS as function of time) of TCO under compressive bending for longer durations of time in open atmosphere shown in Fig. 6d establish the long-term optoelectronic stability of TCO under mechanical flexing. The thermal and chemical stabilities of TCO/ mica are also crucial for optoelectronic device applications. The TCO/mica maintains its inherent metallic character even in bent condition at high temperatures (300 °C) as shown in Fig. 6e unlike those on polymer substrates which curl at temperatures above 150 °C [129] leading to device failure. Furthermore, the TCO/mica electrodes immersed in chemicals like acetone, ethanol, xylene, toluene and chlorobenzene for at least 10 min duration show almost no visible damage on their surface or resistivity change ascertaining their chemical inertness (Fig. 6f). Therefore, epitaxial TCO/mica templates with excellent OptoElectro-Mechanical properties that are robust against chemical and thermal constraints provide a platform to integrate functional devices epitaxially on them and hence, their strong commercialization in flexitronics is highly anticipated. VO2 Vanadium dioxide (VO2) has been seriously investigated as a promising material to implement oxide electronics [130–132] as it undergoes a metal (M)-insulator (I) transition in the vicinity of room temperature. This reversible first-order MI transition is also accompanied by a simultaneous rutile (D4h)-monoclinic (P21/c) structural transformation [132–134] and is strongly coupled to the optical transmittance [130–133,135]. The ultrafast structural transition [136,137] can be induced by thermal [138], electrical [132,135], optical [136] and strain [132,133] stimuli. A flexible

33

VO2 /mica (Fig. 7a) heterostructure is very promising for flexitronics application. Fig. 7b shows the TEM image with an abrupt VO2 /mica interface along the [100] muscovite projection and the corresponding electron diffraction patterns. In between the sharp (00l)-type muscovite reflection spots, the typical [102] VO2 reflections can be seen while the inset shows the electron diffraction pattern of the VO2 film only with the out-of-plane vector corresponding to [020] and the polycrystalline ring of the Pt protection layer deposited during the preparation of the TEM specimen. The epitaxial relationships of (010)VO2k(001)Mica and [100]VO2k[010]Mica consistent with XRD result are obtained. The temperature dependent Raman spectrum in Fig. 7c highlights the structural transition in the VO2/mica heterostructure. Ag and Bg modes [139–144] in the bottom of Fig. 7c indicate that the VO2 at room temperature belongs to an insulating monoclinic phase [139,142] consistent with the XRD and TEM results. As the temperature is raised, the phonon intensities decrease and disappear gradually at around 65 °C implying the monoclinic-rutile structural transformation [138,139]. Moreover, the heavily softened and damped phonons of the rutile phase [139,143] reappear as the temperature is lowered back to room temperature (top of Fig. 7c) confirming the reversible nature of MI transition. The temperature dependent electrical resistance of the VO2/mica in Fig. 7d exhibits a thermal hysteresis (a
9 °C width of the transition region) with 3 orders of resistance change (DR/R > 103) around the Tc (70 °C) characteristic of MI transition. Also, the temperature dependent transmittance at 2600 nm in Fig. 7d through the heating and cooling cycles mimics the thermal hysteresis of the transport data with intensities dropping from 70% to 20% (DTr > 50%) in the vicinity of MI transition. The electrical resistance remains constant (Fig. 7e) under a bending radius down to 5 mm, suggesting the mechanical robustness of this heterostructure. More so, electrical resistance remained unaltered for >103 bending cycles (Fig. 7f). These

Fig. 7. Structure and properties of VO2/mica. (a) Flexible and transparent VO2/mica. (b) Cross-sectional TEM image along with electron diffraction patterns across the VO2/muscovite interface. Temperature dependent (c) Raman spectra and (d) transmittance at 2600 nm and resistance. Resistance as a function of (e) bending radius and (f) bending cycles. Reprinted with permission from Ref. [29]. Copyright 2016 American Chemical Society.

34

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

Fig. 8. Structure and properties of MoO2/mica. (a) Flexible MoO2/mica. (b) Cross-sectional TEM with a sharp interface and corresponding electron diffraction patterns. Temperature dependent (c) Resistance and (d) mobility and carrier concentration. Resistance as a function of (e) bending radius and (f) bending cycles. Reprinted with permission from Ref. [30]. Copyright 2016 American Institute of Physics.

correlated structural-electrical-optical responses of VO2/mica in the vicinity of room temperature with stable performance under mechanical constraints offers a great flexitronics promise especially in the applications of thermal switch and smart window.

MoO2 A vdW epitaxy of metallic MoO2 on muscovite can serve as template for realizing direct heteroepitaxy of other functional oxides necessary for future consumer flexitronics. Fig. 8a shows that MoO2/mica sample is semi-transparent, flexible and can be bent without any observable cracks. Fig. 8b shows a high-resolution (HR) TEM image of a clear and sharp interface between the MoO2 and the muscovite along with the FFT diffraction patterns in the insets. Reciprocal lattices of both MoO2 and muscovite are clearly visible and the epitaxial relationship (010)MoO2||(001)mica and [001]MoO2||[100]mica consistent with the reflection high-energy electron diffraction (RHEED) and XRD results are observed. The bulk-like d-spacing of MoO2 (002) (
2.45 Å) suggests the realization of unconstrained film. More so, a defect-free and non-coherent interface implies the presence of a weak interaction between two materials validating the vdW epitaxy in the MoO2/muscovite heterostructure. The thermal dependence of resistivity (q-T) of the MoO2 in Fig. 8c shows metallic behavior like in the bulk. The MoO2 films are highly conducting (q < 7  104 X cm) at room temperature (300 K) in agreement with those on the conventional substrates [145]. The positive slope of the Hall resistance indicates that the charge carriers in MoO2 are p-type. The temperature evolution of carrier concentration and mobility are shown in Fig. 8d. The room temperature carrier concentration, n, reaches a value of 1.9  1022 cm3, higher than other metal oxides and is almost temperature-independent while the mobility is 0.53 cm2 V1 s1 at 300 K and 0.97 cm2 V1 s1 at 2 K. Absence of magnetic hysteresis and negligible x-ray magnetic circular dichroism signal rules out

the possibility of defect-induced ferromagnetism [146,147] in the high-vacuum deposited MoO2 film. The magnetoresistance MR, MR(%) = 100 X [R(H)  R(0)]/R(0), where R(H) and R(0) are the resistance values with and without magnetic field, respectively, at 2 K is very small (1%) normally expected in metals. The resistance of the MoO2/mica film (<300 nm) under different bending conditions as shown in Fig. 8e remains constant regardless of the bending radius down to 5 mm. The change in resistance after 500 bending cycles under a bending radius of 10 mm shown in Fig. 8f is less than 5%, which asserts the mechanical stability of this heterostructure. Therefore, flexible and semi-transparent MoO2/mica conducting template via vdW epitaxy for flexitronics is realized.

Spintronics (Fe3O4) Spintronics, spin-dependent carrier transport, has opened a new pathway to enhance functionalities of current electronic devices. Fe3O4 is the most attractive material for such applications due to its high Curie temperature (858 K) [148] and is predicted to have nearly 100% spin polarization at the Fermi level [149]. Fe3O4 possesses mixed valent Fe ions (Fe2+, d6; Fe3+, d5) on the octahedral sites and conduction is mediated via the electron hopping between these magnetic ions [150] making it attractive for the magnetic random access memory and magnetic read head applications. To advance the field of flexible spintronics, it is necessary to develop flexible spintronic devices with high-quality single crystalline Fe3O4 films (Fig. 9a). The cross-sectional TEM image represented in Fig. 9b shows a defect-free and sharp Fe3O4/muscovite interface of the heterostructure. The FFT patterns of Fe3O4 and muscovite reveal the epitaxial relationships as Fe3O4(111)||mica(001) and  consistent with XRD and RHEED results. Fe3O4[110]||mica[010], Fig. 9c shows the field-cooled in-plane (IP) magnetization at 100 Oe and zero-field resistance as a function of temperature. The value

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

35

Fig. 9. Structure and properties of Fe3O4/mica. (a) Flexible Fe3O4/mica. (b) Cross-sectional TEM with a sharp interface and corresponding electron diffraction patterns. Temperature dependent (c) Resistance and magnetization exhibiting characteristic Verwey transition feature. Inset shows anisotropy in magnetic hysteresis. (d) Room temperature in-plane magnetic hysteresis at various bending radii. (e) Resistance and (f) magnetoresistance at various radius of bending. Reprinted with permission from Ref. [31]. Copyright 2016 American Chemical Society.

of magnetization doubles from 90 emu/cm3 to 186 emu/cm3 as the temperature increases from 104 K to 132 K, suggesting a magnetic transition temperature at
120 K (the Verwey transition of Fe3O4) which also gets reflected in the transport data. Inset shows anisotropy in magnetic hysteresis at room temperature implying that the magnetic easy axis is along the in-plane direction with the saturated magnetization (Ms) of
410 emu/cm3 and the coercive field (Hc) of
400 Oe. Under a magnetic field of 1 T, the transverse MR values of 0.8%, 1.9%, and 2.1% at 300, 150, and 100 K, respectively comparable to those on rigid substrates [151–153], suggests the feasibility of Fe3O4/muscovite heterostructure in spintronics. The out-of-plane (OOP) magnetization depicted in Fig. 9d shows that the magnetic properties such as the saturation, remanence, and coercivity retain their pristine states while the magnetic anisotropy varies with bending (flex-in mode). The original easy axis of magnetization along the in-plane direction (inset of 9c) gradually rotates to the out-of-plane direction upon bending with a decrease in the magnetic anisotropy. When the bending radius reaches a value of 2.5 mm, the heterostructure becomes isotropic but the initial anisotropic state recovers upon release (see the inset), suggesting tunable magnetic anisotropy of the heterostructure via bending. Furthermore, magnetic force microscopy study also revealed the magnetic domain size increase under bending and recovery to initial domain state after the release. The domain size in a magnetic system is typically determined by several energy terms, including the magnetostatic, magnetocrystalline, magnetostriction, and domain wall energies. Under a compressive strain along the in-plane direction, the easy axis of the positive magnetostrictive [154] Fe3O4 gradually rotates to the out-of-plane direction, resulting in the enlargement of magnetic domain. The macroscopic transport and magnetoresistance properties of the heterostructure under various bending radii are presented in Fig. 9e and Fig. 9f, respectively. The temperature dependent resistance slightly increases under bending but reverts back to the initial state after the strain release however the Verwey transition temperature remains unaltered. The MR values under bent condi-

tions revealed no significant change (around 0.8% under 1 T). The room temperature resistance and MR as a function of bending radius shown in the inset of Fig. 9e display a constant variation. The resistance and MR variation as a function of bending cycles under a bending radius of 5 mm presented in the inset of Fig. 9f show robust performance even after 103 cycles. The verification of the cyclability and endurability of Fe3O4/muscovite heterostructure suggests a new pathway towards flexible spintronic applications.

Magnetostrictive CoFe2O4 Inverse spinel CoFe2O4 (CFO) possesses high curie temperature (793 K) [155], large magnetocrystalline and magnetostrictive anisotropy [156–159]. Conceptually, a CFO/mica biomorph could provide more competitive benefits than that on rigid substrates for optimal performance of flexible magnetostrictive devices. In order to exploit the full potentiality of the flexible magnetostrictive bimorph, the epitaxial growth of CFO on mica is necessary. The theoretical in-plane lattice mismatch between CFO (220) and mica (010) is extremely large (
35%), suggesting a noncoherent growth with a rough interface governed by misfit dislocations. However, experimentally the interface between CFO and mica is very sharp with an ultra thin interdiffused layer (<1 unit cell along c-axis of Mica) without detectable planar defects and grain boundaries as evident from HRTEM image in Fig. 10a. This interdiffused layer has an ambiguous atomic distribution with the mixture of CFO and mica lattices and serves as a seeding layer for the consequent coherent growth. All these salient features are typical of vdW epitaxy discussed earlier. This observation is also supported by irregular RHEED oscillations in the initial stage of the growth, implying a dramatic change of the interface structure between CFO and mica substrate. The FFT patterns from the selected areas of the CFO film and mica substrate indexed in the insets Fig. 10a, confirm the epitaxial relations CFO(111)||mica

36

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

Fig. 10. Structure and properties of CoFe2O4/mica. (a) Cross-sectional TEM with electron diffraction patterns with sharp CoFe2O4/mica interface. (b) Raman spectra of the CFO thin films grown on mica and STO. (c) Raman spectra of CFO thin film under different inward radius of bending. (d) OOP and IP strains as a function of curvature. Magnetic hysteresis loops (e) with the magnetic field along out-of-plane (red) and in-plane (blue) directions, and (f) sample bent along long-axis (orange) and short-axis (olive) with respect to unbent (blue) state. Reprinted with permission from Ref. [32]. Copyright 2017 American Chemical Society.

 (001) and CFO[110]||mica[010] consistent with the RHEED and XRD results. Fig. 10b compares the six Raman modes for CFO on mica and STO(111) reference out of which only five Raman modes: 693 cm1 (A1g), 566 cm1 (T2g(1)), 471 cm1 (T2g(2)), 309 cm1 (Eg) and 183 cm1 (T2g(3)) are active. These modes are unstable against strain and defects leading to cation redistribution [160–162] or short-range ordering [163] resulting in the appearance of an extra mode at
626 cm1. It is clear from the intensity variation and frequency shift of phonons for CFO on mica and STO(111) that CFO has a slight structural distortion. The blue shift in the A1g(1) mode related to the stretching vibrations of tetrahedral coordination along {1 1 1} [161] for the film bent inwards, clearly indicates the decrease in the d-spacing of CFO(111) (Fig. 10c). The lateral bend  ing strain can be expressed as IP ¼  yþt , where y is the distance R from neutral plane inside mica to CFO/mica interface, t is the CFO thickness and R is the radius of bending while the vertical strain can be unveiled from the relation between the wave number and d-spacing of CFO(111) [164] as x0 d0 ¼ xd where x0(x) and d0(d) are the wavenumber and d-spacing of CFO(111) in unbent (bent) state, respectively. The estimated strains as a function of bending curvature in Fig. 10d shows that both OOP and IP CFO lattices suffer the compressive strain when bent inwards such that the OOP strain increases exponentially while the IP strain increases linearly. As shown in Fig. 10e, the magnetic hysteresis loops along the OOP and IP directions are isotropic with a saturation magnetization (Ms) of
150 emu/cm3. The estimated effective anisotropy constants [165] along the OOP and IP directions are 3=2

3=2

6 6 K OOP erg/cm3 and K IP erg/cm3, respeceff
4.31  10 eff
4.44  10 tively, higher than the corresponding single crystal projections along [111] (1.15  106 erg/cm3) and [110] (1.414  106 erg/cm3). This enhancement in the anisotropy constants is ascribed to the slightly distorted CFO lattice as evident from the Raman measure-

ment. A small difference between K OOP and K IP eff eff , imply that the magnetocrystalline anisotropy still dominates the magnetic behavior of the CFO film. Moreover, the room temperature IP magnetization as a function of magnetic field remained unaltered, as shown in Fig. 10f, when the biomorph is compressively bent along either long edge(parallel to CFO[1–10]) or short edge(parallel to CFO[-112]) under a bending radius of 2.5 mm. The MFM images revealed no obvious domain size change independent of the nature of bending strain. The controllable deformation of magnetostrictive CFO/ mica system by magnetic field were investigated via the longitudinal (k||) and transverse (k\) magnetostrictive coefficients with the aid of digital holographic microscopy. The calculated value of k1-10 of ideal single crystal form is 57.2  106, which is in good agreement with the experimental value (k||) of 64  106. However, the calculated value of k11-2 is 117.5  106, significantly differs from the experimental value (k\) of 40  106 with the result that Dk of the CFO film is smaller than that of the ideal CFO single crystal (175 ppm). Such a discrepancy between experimental and theoretical values can be attributed to two possible reasons. First, as the sample is fabricated in a cantilever form so that the preference of magnetization is along the longitudinal edge of sample ([110] direction of this CFO film) due to the shape induced demagnetization state. Additionally, because of the instrumental limitation, the maximum magnetic field used (2 kOe) during the measurement, far from the saturation magnetic field (
30 kOe), led to an underestimation of Dk in the CFO film. Based on these results, the magnetostrictive properties of this flexible CFO/Mica bimorph have extremely high potential for practical applications.

Non-volatile memory (PZT) The demand for high-density data storage, high endurance, and low-power consumption led to the design and development of

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

37

Fig. 11. Structure and properties of PZT/mica. (a) Flexible nonvolatile memory based on PZT/mica. (b) Cross-sectional TEM image depicting the PZT/SRO, SRO/CFO/mica interfaces along with the selected area diffraction patterns of PZT, SRO and mica. (c) out-of-plane PFM phase image. (d) P-E hysteresis loops at various temperatures. (e) Remnant, saturation polarization and coercive field as a function of temperature. (f) Retention and (g) Fatigue measurements at two typical temperatures. Reprinted with permission from Ref. [33]. Copyright 2017 American Association for the Advancement of Science.

flexible non-volatile memory (NVM) elements for next-generation smart electronics, robotics, automotive, healthcare, industrial and military systems. Perovskite Pb(Zr,Ti)O3 (PZT)[166] possessing large polarization, fast polarization switching, high Curie temperature, low coercive field and high piezoelectric coefficient is the most explored system for such applications. Fig. 11a shows a flexible NVM element based on PZT(PbZr0.2Ti0.8O3) on mica. Fig. 11b presents cross-sectional TEM images taken along the zone axis of [010]Mica revealing PZT/SRO, SRO/ CFO/mica interfaces together with the selected area diffraction patterns of PZT, SRO and mica. The sharp interfaces without observable interdiffusion of species across the interfaces indicates high quality of the heterostructure. The epitaxial relationships: (111)SRO||(111)PZT||(001)Mica and [1–10]SRO||[1–10]PZT||[010]Mica consistent with the XRD results are found from the indexed reciprocal lattices. Moreover, the surface morphology probed by atomic force microscope exhibits an average surface roughness (Ra) of 0.62 nm and root-mean-square roughness (Rrms) of 0.81 nm indicating a very smooth surface. In the piezoresponse force microscopy (PFM) measurement, the box-in-box switched patterns were written on the PZT layer via a conducting tip bias of -8 V applied to pole the 3 mm X 3 mm square region followed by another poling with + 8 V in the central area of 1 mm X 1 mm. In Fig. 11c, PFM phase image shows clear bright and dark contrast regions corresponding to upward and downward polarizations, respectively, implying the switchable polarization of PZT layer and hence, the ferroelectric nature of the heterostructure. Well-saturated and symmetric polarization-electric field (P-E) hysteresis loops of the heterostructure measured at 1 MHz and temperatures ranging from 25 °C to 175 °C for a virgin device are shown in Fig. 11d. The temperature evolution of extracted saturation polarization (Psat), remnant polarization (Pr) and coercive field (Ec) shown in Fig. 11e marks the thermal stability of this ferroelectric capacitor. These NVM cells exhibit excellent retention as shown in Fig. 11f even after 105 s at room temperature as well as at high temperatures, best among the flexible NVM elements [167–180]. As displayed in Fig. 11g the PZT/mica capacitor maintained polarization

switching ability at room temperature up to 107 cycles while at high temperature fatigue resistance was maintained until 1010 cycles. It is clear from the aforementioned results that PZT/mica heterostructure can play a key role in high-temperature electronic applications. Fig. 12a tests the P-E hysteresis loops of the PZT capacitors under various compressive and tensile strains while Fig. 12b depicts their performance under tensile and compressive bending of 5 mm before and after 10–1000 bending cycles with no noticeable change. Constant Psat, Pr and Ec, within experimental errors, as a function of bending radius and corresponding strain suggests that the ferroelectric properties of PZT thin film capacitor are stable against mechanical constraints. A sequence of fast voltage pulses for the standard Positive Up and Negative Down (PUND) method [181] was employed to study the polarization switching kinetics. In this method, the device is initially preset to a negative polarized state, P1, at -4 V followed by voltage pulses to measure switching polarization, P2, from negative to positive and switched polarization at 1-4V, P3, in the positive region, respectively. The pulsed switched polarization, DP = P3-P2, as a function of pulse width under various bent conditions is shown in Fig. 12d. The DP (105 mC/cm2) is approximately 1.8 times that of Pr confirming the presence of intrinsic ferroelectricity. The switching speed increases with switching pulse and is unaltered irrespective of the nature of mechanical bending. Fig. 12e displays constant polarization retention of capacitors independent of the type of bending states. It is noteworthy that these capacitors continue to exhibit such robust retention capability even after mechanically flexed for 1000 cycles. Furthermore, Fig. 12f shows stable polarization fatigue behavior in bent states, and remains identical before and after 1000 tensile as well as compressive bending cycles. Polarization remains constant until 107 switching cycles but it reduces to 33% of its initial value after 1010 cycles for the unbent and tensile flexing samples while compressive flexing improves the behavior. Therefore, it is evident from Fig. 12 that the fabricated PZT memory elements exhibit stable and superior performance against mechanical bending highly desirable for flexible device applications.

38

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

Fig. 12. Flexible memory. (a) P-E loops under various bending conditions. (b) P–E hysteresis loops under tensile and compressive bending of 5 mm before and after 10–1000 bending cycles. (c) Psat, Pr and Ec variation as a function of bending radius. (d) Positive-up negative-down (PUND) switching polarization as a function of pulse width at different voltages under mechanical flexing. (e) Retention and (f) Fatigue for the samples in unbent, compressively and tensilely bent before and after 1000 bending cycles. Reprinted with permission from Ref. [33]. Copyright 2017 American Association for the Advancement of Science.

Multiferroics (Self-assembled BFO-CFO nanocomposite) Biphasic magnetoelectric (ME) nanocomposite thin films combining superior electrical and magnetic properties of the constituents have attracted great attention due to their potential memory and sensor applications [182–186]. A combination of multiferroic BiFeO3 (BFO) and ferrimagnetic CFO forms a perfect ME composite due to their large magnetostrictive/piezoelectric coefficients, good lattice match and similar B-site ions mediating interactions across the interfaces. To optimize ME effect, the intimate elastic coupling between the magnetostrictive and piezoelectric materials mediating via the interfaces was tailored adopting advanced thin film architectures like a self-assembled pillarsmatrix (1–3) [182,187,188], quasi-particles-matrix (0–3) [183] and the well-ordered (0–0) nanodots [184,189] structures. However, the substrate clamping effect remains a bottleneck in the realization of ultimate ME coupling. To this end, an alternative strategy of fabricating self-assembled BFO-CFO bulk heterojunction on flexible muscovite via the vdW epitaxy was taken up. The phase-field simulations based on simple energetic arguments exhibits 100% enhancement in the longitudinal ME coupling coefficient, a33, at low magnetic field bias (<500 kA/m) upon the removal of substrate clamping effect as illustrated in Fig. 13a. The photograph of flexible BFO-CFO/mica can be seen in the inset. The opposite IP strains of BFO and CFO obtained from the reciprocal space map points to a strong BFO-CFO structural coupling while surface topography examined by atomic force microscopy shows triangular BFO nanopillars embedded in CFO matrix. The crosssectional TEM along [010] zone axis of the mica depicted in Fig. 13b shows the BFO nanopillars embedded in the CFO matrix with SRO bottom electrode and CFO buffer layer on the mica substrate. HRTEM images show a well-defined phase separation as well as a strong structural coupling among the BFO-CFO phases with sharp interfaces between BFO-CFO/SRO and CFO/mica. The indexed FFT patterns indicate the [100]mica||[111]SRO||[111]CFO||

[111]BFO epitaxial relationships in agreement with the XRD result. Furthermore, isotropic magnetic hysteresis loops along the IP and OOP directions with the saturated magnetization (Ms)
237 emu/ cm3 and the coercivity (Hc)
2 kOe were observed. Also, M-H loops under the different bending modes (bending radius
2 mm) remain unaltered as shown in Fig. 13c similar to CFO/mica (Fig. 10f). The OOP converse piezoelectric coefficient, d33, of BFOCFO on mica (76.5 pmV1) higher than that on STO(111) (50 pmV1) grown under identical conditions highlights the role of the substrate constraint release in enhancing piezoelectric response of the BFO pillars. The d33 of each selected individuals BFO pillars under various tensile and compressive bends at H = 0, 0.1, 0.2 T are depicted in Fig. 13d. For the magnetic field applied in OOP direction, d33 decreases due to the negative longitudinal magnetostriction of CFO matrix which compresses BFO pillars via their strong structural coupling. Under a bending strain, d33, reduces further due to the variation of length of BFO pillars along c-axis and hence, the rotation of P. The associated magnetoelectric coefficient (aME) is estimated by aME = DE/DH where DE = Du/ (d33D), Du is the piezo-response displacement with and without magnetic bias, D is the thickness of BFO-CFO nanocomposite (
200 nm), and the corresponding d33 without magnetic bias (76.5 pm/V) [183,184,189]. The calculated aME of BFO-CFO/mica (
11,000 mV/cmOe) is much higher than that found in the quasi (0–3) BFO-CFO/STO (338 mV/cmOe) [183] and comparable to unclamped PZT/CFO core–shell nanofiber structure (29,500 mV/ cmOe) [189]. The variation of local electric response of the BFO pillars under mechanical strain dictates the variation of aME. However, the aME varies much slower than local d33 under mechanical strain (Fig. 13d and e). Notably, the room temperature macroscopic magnetoelectric susceptibility of BFO-CFO/mica is about
74 mV/cmOe (24% enhancement) higher than BFO-CFO/ STO(001) (
60 mV/cmOe) [182] displaying similar (1–3) nanostructure in accordance with the phase-field simulations. This value is best among the ME systems reported on rigid as well as flexible

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

39

Fig. 13. (a) Longitudinal magnetoelectric coupling coefficient, a33, simulated using a phase-field model, which shows an enhancement of over 100% at low magnetic field bias (<500 kA/m) on removing the substrate clamping. (b) Cross-sectional TEM image of the BFO-CFO/SRO/CFO/mica heterostructure. HRTEM images of the BFO-CFO/SRO and CFO/mica interfaces along with indexed FFT patterns are shown. (c) Room temperature M-H loops under bending modes. (d) d33 of each selected individual BFO pillars under different bending radii with and without applied magnetic field. (e) Calculated aME from (d). (f) aME variation as a function of bending cycles. Reprinted with permission from Ref. [34]. Copyright 2017 American Chemical Society.

substrates [190–196]. The aME variation as a function of bending cycles under a bending radius of 7 mm shown in Fig. 13f exhibits robust values up to 1000 cycles indicating its mechanical stability. Thus, BFO-CFO/mica with excellent structural quality and unprecedented ME coupling against mechanical strain provides a new direction in the realization of highly sensitive yet flexible ME electronic devices application with optimized performance. Outlook and future scope Although historically developed for planar structures, vdW epitaxy can be a prospective route towards successful heteroepitaxial growth of different geometries and device architectures. More interesting geometries like vdW laminates [1], stacking of distinct 2D layered systems exhibiting magnetism, ferroelectricity, multiferroicity, superconductivity and so on, studying proximity effects can offer unprecedented opportunities to explore novel quantum phenomena and realize new applications. From the ongoing critical assessment of vdW epitaxy on mica, it is evident that the MICAtronics will find potential applications in the field of flexible electronics, realize exotic 2D phases at oxide interface and tailor new functionalities. Furthermore, it would open a new window for creating new type of flexible device applications as they do not involve unpractical and tedious transfer processes. Mica based thin films and heterointerfaces serve as model systems to study fundamental nucleation and growth processes in vdW epitaxy and the different mechanisms that govern the physical properties of unconstrained heterostructures. The most valuable advantage is that these vdW heterodevices are resilient to extreme conditions (thermal, mechanical and chemical). This is in stark contrast with the rough surfaces and high-temperature inability of the conventional flexible soft substrates that have detrimental effect on device performance. Furthermore, compatibility of mica with cur-

rent Si based technologies can open up promising avenue for future microelectronics. Moreover, our process can be transferred to the sputtering process for large scale production. Thus, very complex nanometer scale device engineering for optimized device performance can be envisaged combined with the aforementioned benefits. Due to shapeability and light-weight, device miniaturization demanded by the industry, MICAtronics can take a new leap. For example, reducing the size as well as making flexible devices will have greater impact in biomedical industry like miniaturized heart implants unlike the currently used bulky position measurement modules. Moreover, by designing compact flexible device elements with reduced power consumption capability will serve as big leap in energy savings in heavy instruments/machines. Therefore, highperformance, cheap and eco-friendly technologies can be envisioned via MICAtronics. Conclusions In conclusion, high quality oxide heteroepitaxy on flexible and transparent muscovite for future flexible device applications has been successfully demonstrated. Furthermore, these systems retain superior performance of their epitaxial counterparts on rigid substrates and surpass those on soft flexible substrates. Our process can be transferred to the sputtering, CVD methods and solution-processed techniques (inkjet printing, screen printing and spray coating) for large scale production. Thus, MICAtronics can undeniably open up a new exciting avenue to integrate functional oxides for next generation multifunctional devices that are robust against thermal, mechanical and chemical constraints. Additionally, the development of new material designs and architectures for desired applications with scaling up capability can catch up soon with competing mature flexitronics technologies. There awaits numerous exciting opportunities in developing the

40

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

growth of high-quality materials on muscovite mica which would not only expand our understanding of the underlying physics but also potentially lead to the discovery of unanticipated phenomena and applications.

Acknowledgments This work was supported by the Ministry of Science and Technology – Taiwan (ROC).

References [1] A.K. Geim, I.V. Grigorieva, Nature 499 (2014) 419–425. [2] R. Mas-Balleste, C. Gómez-Navarro, J. Gómez-Herrero, F. Zamora, Nanoscale 3 (2011) 20–30. [3] M. Xu, T. Liang, M. Shi, H. Chen, Chem. Rev. 113 (2013) 3766–3798. [4] S.Z. Butler, S.M. Hollen, L. Cao, Y. Cui, J.A. Gupta, H.R. Gutierrez, T.F. Heinz, S.S. Hong, J. Huang, A.F. Ismach, E. Johnston-Halperin, M. Kuno, V.V. Plashnitsa, R. D. Robinson, R.S. Ruoff, S. Salahuddin, J. Shan, L. Shi, M.G. Spencer, M. Terrones, W. Windl, J.E. Goldberger, ACS Nano 7 (2013) 2898–2926. [5] K.S. Novoselov, A. Mishchenko, A. Carvalho, A.H. Castro Neto, Science 353 (2016) aac9439. [6] D. Jariwala, T.J. Marks, M.C. Hersam, Nat. Mater. 16 (2017) 170–181. [7] S. Das, J.A. Robinson, M. Dubey, H. Terrones, M. Terrones, Annu. Rev. Mater. Res. 45 (2015) 1–27. [8] C.H. Lee, T. Schiros, E.J. Santos, B. Kim, K.G. Yager, S.J. Kang, S. Lee, J. Yu, K. Wantanbe, T. Taniguchi, J. Hone, E. Kaxiras, C. Nuckolls, P. Kim, Adv. Mater. 26 (2014) 2812–2817. [9] G.H. Lee, C.H. Lee, A.M. van der Zande, M. Han, X. Cui, G. Arefe, C. Nuckolls, T.F. Heinz, J. Hone, P. Kim, APL Mater. 2 (2014) 092511. [10] O. Lopez-Sanchez, E. Alarcon Llado, V. Koman, A. Fontcuberta i Morral, A. Radenovic, A. Kis, ACS Nano 8 (2014) 3042–3048. [11] F. Withers, O.D. Pozo-Zamudio, A. Mishchenko, A.P. Rooney, A. Gholinia, K. Watanabe, T. Taniguchi, S.J. Haigh, A.K. Geim, A.I. Tartakovskii, K.S. Novoselov, Nat. Mater. 14 (2015) 301–306. [12] A.C. Ferrari, F. Bonaccorso, V. Falko, K.S. Novoselov, S. Roche, P. Bøggild, S. Borini, F. Koppens, V. Palermo, N. Pugno, J.A. Garrido, R. Sordan, A. Bianco, L. Ballerini, M. Prato, E. Lidorikis, J. Kivioja, C. Marinelli, T. Ryhänen, A. Morpurgo, J.N. Coleman, V. Nicolosi, L. Colombo, A. Fert, M. GarciaHernandez, A. Bachtold, G.F. Schneider, F. Guinea, C. Dekker, M. Barbone, C. Galiotis, A. Grigorenko, G. Konstantatos, A. Kis, M. Katsnelson, C.W.J. Beenakker, L. Vandersypen, A. Loiseau, V. Morandi, D. Neumaier, E. Treossi, V. Pellegrini, M. Polini, A. Tredicucci, G.M. Williams, B.H. Hong, J.H. Ahn, J.M. Kim, H. Zirath, B.J. van Wees, H. van der Zant, L. Occhipinti, A. Di Matteo, I.A. Kinloch, T. Seyller, E. Quesnel, X. Feng, K. Teo, N. Rupesinghe, P. Hakonen, S.R. T. Neil, Q. Tannock, T. Löfwander, J. Kinaret, Nanoscale 7 (2014) 4598–4810. [13] S.M. Auerbach, K.A. Carrado, P.K. Dutta (Eds.), Handbook of Layered Materials, Marcel Dekker, New York, 2004. [14] S. Lin, Y. Chui, Y. Li, S.P. Lau, FlatChem 2 (2017) 15–37. [15] Y. Celik, E. Flahaut, E. Suvaci, FlatChem 1 (2017) 74–88. [16] E.O. Ortiz-Quiles, C.R. Cabrera, FlatChem 2 (2017) 1–7. [17] L. Rounan, Z. Peinan, J. Non-Cryst, Solids 84 (1986) 600–604. [18] J.M. Tzeng, J.G. Duh, K.H. Chung, C.C. Chan, J. Mater. Sci. 28 (1992) 6127–6135. [19] L. Cheng, P. Fenter, K.L. Nagy, M.L. Schlegel, N.C. Sturchio, Phys. Rev. Lett. 87 (2001) 156103. [20] T. Fukuma, Y. Ueda, S. Yoshioka, H. Asakawa, Phys. Rev. Lett. 104 (2010) 016101. [21] H. Bluhm, T. Inoue, M. Salmeron, Surf. Sci. 462 (2000) L599–L601. [22] P.B. Miranda, L. Xu, Y.R. Shen, M. Salmeron, Phys. Rev. Lett. 81 (1998) 5876– 5879. [23] C. Spagnoli, K. Loos, A. Ulman, M.K. Cowman, J. Am. Chem. Soc. 125 (2003) 7124–7128. [24] H.K. Christenson, N.H. Thomson, Surf. Sci. Rep. 71 (2016) 367–390. [25] Y. He, H. Dong, Q. Meng, L. Jiang, W. Shao, L. He, W. Hu, Adv. Mater. 23 (2011) 5502–5507. [26] C.G. Low, Q. Zhang, Small 8 (2012) 2178–2183. [27] X. Zhang, Y. He, R. Li, H. Dong, W. Hu, Adv. Mater. 28 (2016) 3755–3760. [28] Y. Bitla, C. Chen, H.C. Lee, T.H. Do, C.H. Ma, L.V. Qui, C.W. Huang, W.W. Wu, L. Chang, P.W. Chiu, Y.H. Chu, A.C.S. Appl, Mater. Interfaces 8 (2016) 32401– 32407. [29] C.I. Li, J.C. Lin, H.J. Liu, M.W. Chu, H.W. Chen, C.H. Ma, C.Y. Tsai, H.W. Huang, H. J. Lin, H.L. Liu, P.W. Chiu, Y.H. Chu, Chem. Mater. 28 (2016) 3914–3919. [30] C.H. Ma, J.C. Lin, H.J. Liu, T.H. Do, Y.M. Zhu, T.D. Ha, Q. Zhan, J.Y. Juang, Q. He, E. Arenholz, P.W. Chiu, Y.H. Chu, Appl. Phys. Lett. 108 (2016) 253104. [31] P.C. Wu, P.F. Chen, T.H. Do, Y.H. Hsieh, C.H. Ma, T.D. Ha, K.H. Wu, Y.J. Wang, H. B. Li, Y.C. Chen, J.Y. Juang, P. Yu, L.M. Eng, C.F. Chang, P.W. Chiu, L.H. Tjeng, Y. H. Chu, A.C.S. Appl, Mater. Interfaces 8 (2016) 33794–33801. [32] H.J. Liu, C.K. Wang, D. Su, T. Amrillah, Y.H. Hsieh, K.H. Wu, Y.C. Chen, J.Y. Juang, L.M. Eng, S.U. Jen, Y.H. Chu, A.C.S. Appl, Mater. Interfaces 9 (2017) 7297–7304.

[33] J. Jiang, Y. Bitla, C.W. Huang, T.H. Do, H.J. Liu, Y.H. Hsieh, C.H. Ma, C.Y. Jang, Y. H. Lai, P.W. Chiu, W.W. Wu, Y.C. Chen, Y.C. Zhou, Y.H. Chu, Sci. Adv. 3 (2017) e1700121. [34] T. Amrillah, Y. Bitla, K. Shin, T. Yang, Y.H. Hsieh, Y.Y. Chiou, H.J. Liu, T.H. Do, D. Su, Y.C. Chen, S.U. Jen, L.Q. Chen, K.H. Kim, J.Y. Juang, Y.H. Chu, ACS Nano, http://dx.doi.org/10.1021/acsnano.7b02102. [35] H. Usami, A. Nakasa, M. Adachi, E. Suzuki, H. Fujimatsu, T. Ohashi, S. Yamada, K. Tsugita, Y. Taniguchi, Thin Solid Films 515 (2006) 2310–2315. [36] S. Ke, C. Chen, N. Fu, H. Zhou, M. Ye, P. Lin, W. Yuan, X. Zeng, L. Chen, H. Huang, A.C.S. Appl, Mater. Interfaces 8 (2016) 28406–28411. [37] M.A. Herman, H. Sitter, Molecular Beam Epitaxy, Springer, Berlin, 1996. [38] A. Koma, K. Sunouchi, T. Miyajima, Microelectron. Eng. 2 (1984) 129–136. [39] A. Koma, K. Sunouchi, T. Miyajima, J. Vac. Sci. Technol., B 3 (1985) 724. [40] A. Koma, K. Yoshimura, Surf. Sci. 174 (1986) 556–560. [41] A. Koma, K. Ueno, K. Saiki, J. Cryst. Growth 111 (1991) 1029–1032. [42] A. Koma, Thin Solid Films 216 (1992) 72–76. [43] A. Koma, J. Cryst. Growth 201 (202) (1999) 236–241. [44] S.C. Ying, Phys. Rev. B 3 (1971) 4160–4171. [45] A.D. Novaco, J.P. McTague, Phys. Rev. Lett. 38 (1977) 1286–1289. [46] B.S. Mitchell, An Introduction to Materials Engineering and Science: For Chemical and Materials Engineers, Wiley, New York, 2008. [47] W. Jaegermann, A. Klein, C. Pettenkofer, Electronic Properties of van der Waals-Epitaxy Film and Interfaces, in: H.P. Hughes, H.I. Starnberg (Eds.), Electron Spectroscopies Applied to Low-Dimensional Materials, Springer, Netherlands, 2002, pp. 317–402. [48] M.I.B. Utama, M. de la Mata, C. Magen, J. Arbiol, Q. Xiong, Adv. Funct. Mater. 23 (2013) 1636–1646. [49] L. Pauling, Proc. Natl. Acad. Sci. 16 (1930) 123–129. [50] Ch. Mauguin, Comp. Rend. 185 (1927) 288–291. [51] A. Castellanos-Gomez, M. Poot, A. Amor-Amorós, G.A. Steele, H.S.J. van der Zant, N. Agraït, G. Rubio-Bollinger, Nano Res. 5 (2012) 550–557. [52] S.G. Barlow, D.A.C. Manning, Br. Ceram. Trans. 98 (1999) 122–126. [53] S. Steinberg, W. Ducker, G. Vigil, C. Hyukjin, C. Frank, M.Z. Tseng, D.R. Clarke, J. N. Israelachvili, Science 260 (1993) 656–659. [54] J.T. Kendall, D. Yeo, Proc. Phys. Soc. Sect. B 64 (2002) 135–142. [55] D. McCoul, W. Hu, M. Gao, V. Mehta, Q. Pei, Adv. Electron Mater. 2 (2016) 1500407. [56] A.M. Hussain, M.M. Hussain, Adv. Mater. 28 (2016) 4219–4249. [57] W.S. Wong, A. Sallelo, Flexible Electronics: Materials and Applications, Springer, New York, 2009. [58] L. Nyholm, G. Nystrom, A. Mihranyan, M. Stromme, Adv. Mater. 23 (2011) 3751–3769. [59] D. Tobjork, R. Osterbacka, Adv. Mater. 23 (2011) 1935–1961. [60] L.B. Hu, Y. Cui, Energy Environ. Sci. 5 (2012) 6423–6435. [61] W. Zeng, L. Shu, Q. Li, S. Chen, F. Wang, X.M. Tao, Adv. Mater. 26 (2014) 5310– 5336. [62] K.D. Harris, A.L. Elias, H.J. Chung, J. Mater. Sci. 51 (2016) 2771–2805. [63] Z. Suo, E.Y. Ma, H. Gleskova, S. Wagner, Appl. Phys. Lett. 74 (1999) 1177– 1179. [64] P. Yu, Y.H. Chu, R. Ramesh, Mater. Today 15 (2012) 320–327. [65] L.W. Martin, Y.H. Chu, R. Ramesh, Mater. Sci. Eng., R 68 (2010) 89–133. [66] D.G. Schlom, L.Q. Chen, X. Pan, A. Schmehl, M.A. Zurbuchen, J. Am. Ceram. Soc. 91 (2008) 2429–2454. [67] B. Putsche, L. Tumbek, A. Winkler, J. Chem. Phys. 137 (2012) 134701. [68] G. Lippert, J. Dabrowski, M. Lemme, C. Marcus, O. Seifarth, G. Lupina, Phys. Status Solidi B 248 (2011) 2619–2622. [69] G. Lippert, J. Dabrowski, Y. Yamamoto, F. Herziger, J. Maultzsch, M.C. Lemme, W. Mehr, G. Lupina, Carbon 52 (2013) 40–48. [70] K. Schouteden, B. Amin-Ahmadi, Z. Li, D. Muzychenko, D. Schryvers, C. Van Haesendonck, Nat. Commun. 7 (2016) 14001. [71] K. Schouteden, C. Van Haesendonck, Phys. Rev. Lett. 108 (2012) 076806. [72] H. Tanaka, M. Taniguchi, Jpn. J. Appl. Phys. 55 (2016) 078003. [73] K. Ueno, K. Saiki, T. Shimada, A. Koma, J. Vac. Sci. Technol., A 8 (1990) 68–72. [74] Q. Ji, Y. Zhang, T. Gao, Y. Zhang, D. Ma, M. Liu, Y. Chen, X. Qiao, P.H. Tan, M. Kan, J. Feng, Q. Sun, Z. Liu, Nano Lett. 13 (2013) 3870–3877. [75] H. Peng, W. Dang, J. Cao, Y. Chen, D. Wu, W. Zheng, H. Li, Z.X. Shen, Z. Liu, Nat. Chem. 4 (2012) 281–286. [76] H. Li, J. Cao, W. Zheng, Y. Chen, D. Wu, W. Dang, K. Wang, H. Peng, Z. Liu, J. Am. Chem. Soc. 134 (2012) 6132–6135. [77] N.H. Tu, Y. Tanabe, Y. Satake, K.K. Huynh, K. Tanigaki, Nat. Commun 7 (2016) 13763–13768. [78] D.Y. Kim, S. Lee, Z.H. Lin, K.W. Choi, S.G. Doo, H. Chang, J.Y. Leem, Z.L. Wang, S. O. Kim, Nano Energy 9 (2014) 101–111. [79] M.I.B. Utama, F.J. Belarre, C. Magen, B. Peng, J. Arbiol, Q. Xiong, Nano Lett. 12 (2012) 2146–2152. [80] M.I.B. Utama, Q. Zhang, S. Jia, D. Li, J. Wang, Q. Xiong, ACS Nano 6 (2012) 2281–2288. [81] M. Gillet, R. Delamare, E. Gillet, J. Cryst. Growth 279 (2005) 93–99. [82] M.I.B. Utama, Q. Zhang, J. Zhang, Y. Yuan, F.J. Belarre, J. Arbiol, Q.H. Xiong, Nanoscale 5 (2013) 3570–3588. [83] K. Ellmer, Nat. Photonics 6 (2012) 809–817. [84] D.S. Ginley, H. Hosono, D.C. Paine, Handbook of Transparent Conductors, Springer Series in Materials Science, New York, 2010, pp. 1–264. [85] H. Ohta, M. Orita, M. Hirano, H. Tanji, H. Kawazoe, H. Hosono, Appl. Phys. Lett. 76 (2000) 2740–2742. [86] H. Ohta, M. Orita, M. Hirano, J. Appl. Phys. 91 (2002) 3547–3550.

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42 [87] H. Kim, J.S. Horwitz, W.H. Kim, Z.H. Kafafi, D.B. Chrisey, J. Appl. Phys. 91 (2002) 5371–5376. [88] H. Kim, J.S. Horwitz, S.B. Qadri, D.B. Chrisey, Thin Solid Films 420–421 (2002) 107–111. [89] C.T. Riley, J.S. Smalley, K.W. Post, D.N. Basov, Y. Fainman, D. Wang, Z. Liu, D.J. Sirbuly, Small 12 (2016) 892–901. [90] Q. Wan, E.N. Dattoli, W.Y. Fung, W. Guo, Y. Chen, X. Pan, W. Lu, Nano Lett. 6 (2006) 2909–2915. [91] J.D. Emery, C.M. Schleputz, P.J. Guo, R.P.H. Chang, A.B.F. Martinson, Cryst. Growth Des. 16 (2016) 640–645. [92] J.A. Jeong, H.S. Shin, K.H. Choi, H.K. Kim, J. Phys. D Appl. Phys. 43 (2010) 465403. [93] M. Girtan, M. Kompitsas, R. Mallet, I. Fasaki, Eur. Phys. J. Appl. Phys. 51 (2010) 33212. [94] J.P. Kim, S.A. Lee, S.J. Bae, S.K. Park, U.C. Choi, C.R. Cho, Thin Solid Films 516 (2008) 5223–5226. [95] S.M. Park, T. Ikegami, K. Ebihara, P.K. Shin, Appl. Surf. Sci. 253 (2006) 1522– 1527. [96] Y. Liu, L. Zhao, J.S. Lian, Vacuum 81 (2006) 18–21. [97] Y. Liu, J. Lian, Appl. Surf. Sci. 253 (2007) 3727–3730. [98] C. Agashe, O. Kluth, J. Hüpkes, U. Zastrow, B. Rech, M. Wuttig, J. Appl. Phys. 95 (2004) 1911–1917. [99] T.L. Yang, D.H. Zhang, J. Ma, H.L. Ma, Y. Chen, Thin Solid Films 326 (1998) 60– 62. [100] H.L. Shen, H. Zhang, L.F. Lu, F. Jiang, C. Yang, Prog. Nat. Sci. Mater. Int. 20 (2010) 44–48. [101] H. Agura, A. Suzuki, T. Matsushita, T. Aoki, M. Okuda, Thin Solid Films 445 (2003) 263–267. [102] R.J. Mendelsberg, S.H.N. Lim, Y.K. Zhu, J. Wallig, D.J. Milliron, A. Anders, J. Phys. D Appl. Phys. 44 (2011) 232003. [103] D.J. Lee, H.M. Kim, J.Y. Kwon, H. Choi, S.H. Kim, K.B. Kim, Adv. Funct. Mater. 21 (2011) 448–455. [104] Y. Kim, W. Lee, D.R. Jung, J. Kim, S. Nam, H. Kim, B. Park, Appl. Phys. Lett. 96 (2010) 171902. [105] S. Cornelius, M. Vinnichenko, N. Shevchenko, A. Rogozin, A. Kolitsch, W. Moller, Appl. Phys. Lett. 94 (2009) 042103. [106] H. Sakata, S. Noguchi, J. Phys. D Appl. Phys. 14 (1981) 1523–1529. [107] J.H. Kim, K.A. Jeon, G.H. Kim, S.Y. Lee, Appl. Surf. Sci. 252 (2006) 4834–4837. [108] H.R. Choi, S.K. Eswaran, S.M. Lee, Y.S. Cho, A.C.S. Appl, Mater. Interfaces 7 (2015) 17569–17572. [109] J. Yun, Y.H. Park, T.S. Bae, S. Lee, G.H. Lee, A.C.S. Appl, Mater. Interfaces 5 (2013) 164–172. [110] H. Wu, L. Hu, T. Carney, Z. Ruan, D. Kong, Z. Yu, Y. Yao, J.J. Cha, J. Zhu, S. Fan, Y. Cui, J. Am. Chem. Soc. 133 (2011) 27–29. [111] N. Kim, H. Kang, J.H. Lee, S. Kee, S.H. Lee, K. Lee, Adv. Mater. 27 (2015) 2317– 2323. [112] L.M. Wang, Y.J. Chen, J.W. Liao, J. Phys. Chem. Sol. 69 (2008) 527–530. [113] S.K. Park, J.I. Han, W.K. Kim, M.G. Kwak, Thin Solid Films 397 (2001) 49–55. [114] M.G. Sandoval-Paz, R. Ramirez-Bon, Thin Solid Films 517 (2009) 2596–2601. [115] H. Han, J.W. Mayer, T.L. Alford, J. Appl. Phys. 99 (2006) 123711. [116] C. Yang, S. Lee, T. Lin, S. Chen, Thin Solid Films 516 (2008) 1984–1991. [117] C. Guillen, J. Herrero, Mater. Chem. Phys. 112 (2008) 641–644. [118] D.H. Kim, M.R. Park, G.H. Lee, Surf. Coat. Technol. 201 (2006) 927–931. [119] C. Guillen, J. Herrero, Semicond. Sci. Technol. 23 (2008) 075002. [120] L. Hao, X. Diao, H. Xu, B. Gu, T. Wang, Appl. Surf. Sci. 254 (2008) 3504–3508. [121] J. Lee, H. Jung, J. Lee, D. Lim, K. Yang, J. Yi, W.C. Song, Thin Solid Films 516 (2008) 1634–1639. [122] G. Haacke, J. Appl. Phys. 47 (1976) 4086–4089. [123] S. Ray, R. Banerjee, N. Basu, A. Batabyal, A. Barua, J. Appl. Phys. 54 (1983) 3497–3501. [124] D. Paeng, J.H. Yoo, J. Yeo, D. Lee, E. Kim, S.H. Ko, C. Grigoropoulos, Adv. Mater. 27 (2015) 2762–2767. [125] H.G. Im, B.W. An, J. Jin, J. Jang, Y.G. Park, J.U. Park, B.S. Bae, Nanoscale 8 (2016) 3916–3922. [126] C. Feng, K. Liu, J.S. Wu, L. Liu, J.S. Cheng, Y. Zhang, Y. Sun, Q. Li, S. Fan, K. Jiang, Adv. Funct. Mater. 20 (2010) 885–891. [127] B. Deng, P.C. Hsu, G. Chen, B.N. Chandrashekar, L. Liao, Z. Ayitimuda, J. Wu, Y. Guo, L. Lin, Y. Zhou, M. Aisijiang, Q. Xie, Y. Cui, Z. Liu, H. Peng, Nano Lett. 15 (2015) 4206–4213. [128] Z. Chen, B. Cotterell, W. Wang, E. Guenther, S.J. Chua, Thin Solid Films 394 (2001) 201–205. [129] V. Zardetto, T.M. Brown, A. Reale, A. Di Carlo, J. Polym. Sci. B Polym. Phys. 49 (2011) 638–648. [130] Z. Yang, C. Ko, S. Ramanathan, Annu. Rev. Mater. Res. 41 (2011) 337–367. [131] S. Hormoz, S. Ramanathan, Solid-State Electron. 54 (2010) 654–659. [132] M. Nakano, K. Shibuya, D. Okuyama, T. Hatano, S. Ono, M. Kawasaki, Y. Iwasa, Y. Tokura, Nature 487 (2012) 459–462. [133] N.B. Aetukuri, A.X. Gray, M. Drouard, M. Cossale, L. Gao, A.H. Reid, H. Ohldag, C.A. Jenkins, E. Arenholz, K.P. Roche, H.A. Dürr, M.G. Samant, S.S.P. Parkin, R. Kukreja, Nat. Phys. 9 (2013) 661–666. [134] M. Imada, A. Fujimori, Y. Tokura, Rev. Mod. Phys. 70 (1998) 1039–1236. [135] J. Jeong, N. Aetukuri, T. Graf, T.D. Schladt, M.G. Samant, S.S.P. Parkin, Science 339 (2013) 1402–1405. [136] A. Cavalleri, T. Dekorsy, H.H.W. Chong, J.C. Kieffer, R.W. Schoenlein, Phys. Rev. B 70 (2004) 161102R.

41

[137] H.T. Kim, Y.W. Lee, B.J. Kim, B.G. Chae, S.J. Yun, K.Y. Kang, K.J. Han, K.J. Yee, Y.S. Lim, Phys. Rev. Lett. 97 (2006) 266401. [138] H. Kim, Y. Kim, K.S. Kim, H.Y. Jeong, A.R. Jang, S.H. Han, D.H. Yoon, K.S. Suh, H. S. Shin, T.Y. Kim, W.S. Yang, ACS Nano 7 (2013) 5769–5776. [139] R. Srivastava, L.L. Chase, Phys. Rev. Lett. 27 (1971) 727–730. [140] V.G. Golubev, V.Y. Davydov, N.F. Kartenko, D.A. Kurdyukov, A.V. Medvedev, A. B. Pevtsov, A.V. Scherbakov, E.B. Shadrin, Appl. Phys. Lett. 79 (2001) 2127– 2129. [141] G.I. Petrov, V.V. Yakovlev, J. Squier, Appl. Phys. Lett. 81 (2002) 1023–1025. [142] P. Schilbe, Phys. B 316–317 (2002) 600–602. [143] H.T. Kim, B.G. Chae, D.H. Youn, G. Kim, K.Y. Kang, S.J. Lee, K. Kim, Y.S. Lim, Appl. Phys. Lett. 86 (2005) 242101. [144] L. Bai, Q. Li, S.A. Corr, Y. Meng, C. Park, S.V. Sinogeikin, C. Ko, J. Wu, G. Shen, Phys. Rev. B 91 (2015) 104110. [145] V. Bhosle, A. Tiwari, J. Narayan, J. Appl. Phys. 97 (2005) 083539. [146] P. Thakur, J.C. Cezar, N.B. Brookes, R.J. Choudhary, R. Prakash, D.M. Phase, K.H. Chae, R. Kumar, Appl. Phys. Lett. 94 (2009) 062501. [147] J. Li, S. Zhang, J. Bartell, C. Nisoli, X. Ke, P.E. Lammert, V.H. Crespi, P. Schiffer, Phys. Rev. B 82 (2010) 134407. [148] P.A. Miles, W.B. Westphal, A. Von Hippel, Rev. Mod. Phys. 29 (1957) 279–307. [149] A. Yanase, K. Siratori, J. Phys. Soc. Jpn. 53 (1984) 312–317. [150] A. Rosencwaig, Can. J. Phys. 47 (1969) 2309–2317. [151] M. Ziese, H.J. Blythe, J. Phys.: Condens. Matter 12 (2000) 13–28. [152] S.B. Ogale, K. Ghosh, R.P. Sharma, R.L. Greene, R. Ramesh, T. Venkatesan, Phys. Rev. B 57 (1998) 7823–7828. [153] S.K. Arora, R.G.S. Sofin, I.V. Shvets, Phys. Rev. B 72 (2005) 134404. [154] R.C. O’Handley, Modern Magnetic Materials: Principles and Applications, Wiley, 2000, p. 225. [155] L. Bouet, P. Tailhades, A. Rousset, J. Magn. Magn. Mater. 153 (1996) 389–396. [156] R.M. Bozorth, E.F. Tilden, A.J. Williams, Phys. Rev. 99 (1955) 1788–1798. [157] J.C. Slonczewski, Phys. Rev. 110 (1958) 1341–1348. [158] Y.X. Zheng, Q.Q. Cao, C.L. Zhang, H.C. Xuan, L.Y. Wang, D.H. Wang, Y.W. Du, J. Appl. Phys. 110 (2011) 043908. [159] D. Fritsch, C. Ederer, Phys. Rev. B 82 (2010) 104117. [160] P. Chandramohan, M.P. Srinivasan, S. Velmurugan, S.V. Narasimhan, J. Solid State Chem. 184 (2011) 89–96. [161] Y.Y. Liao, Y.W. Li, Z.G. Hu, J.H. Chu, Appl. Phys. Lett. 100 (2012) 071905. [162] H.J. Liu, Y.Y. Liu, C.Y. Tsai, S.C. Liao, Y.J. Chen, H.J. Lin, C.H. Lai, W.F. Hsieh, J.Y. Li, C.T. Chen, Q. He, Y.H. Chu, Sci. Rep. 5 (2015) 12073. [163] V.G. Ivanov, M.V. Abrashev, M.N. Iliev, M.M. Gospodinov, J. Meen, M.I. Aroyo, Phys. Rev. B 82 (2010) 024104. [164] M. Foerster, M. Iliev, N. Dix, X. Martí, M. Barchuk, F. Sánchez, J. Fontcuberta, Adv. Funct. Mater. 22 (2012) 4344–4351. [165] T. Dhakal, D. Mukherjee, R. Hyde, P. Mukherjee, M.H. Phan, H. Srikanth, S. Witanachchi, J. Appl. Phys. 107 (2010) 053914. [166] N. Izyumskaya, Y.I. Alivov, S.J. Cho, H. Morkoç, H. Lee, Y.S. Kang, Crit. Rev. Solid State Mater. Sci. 32 (2007) 111–202. [167] M.T. Ghoneim, M.A. Zidan, M.Y. Alnassar, A.N. Hanna, J. Kosel, K.N. Salama, M. M. Hussain, Adv. Electron. Mater. 1 (2015) 1500045–1500054. [168] M.T. Ghoneim, M.M. Hussain, Appl. Phys. Lett. 107 (2015) 052904–052908. [169] Z. Zuo, B. Chen, Q.F. Zhan, Y. Liu, H. Yang, Z. Li, G. Xu, R.W. Li, J. Phys. D Appl. Phys. 45 (2012) 185302–185306. [170] A.I. Kingon, S. Srinivasan, Nat. Mater. 4 (2005) 233–237. [171] C.T. Shelton, B.J. Gibbons, J. Am. Ceram. Soc. 94 (2011) 3223–3226. [172] J. Rho, S.J. Kim, W. Heo, N.E. Lee, H.S. Lee, J.H. Ahn, I.E.E.E. Ele. Dev. Lett. 31 (2010) 1017–1019. [173] I. Bretos, R. Jiménez, A. Wu, A.I. Kingon, P.M. Vilarinho, M.L. Calzada, Adv. Mater. 26 (2014) 1405–1409. [174] W.Y. Kim, H.C. Lee, I.E.E.E. Ele, Dev. Lett. 33 (2012) 260–262. [175] D. Mao, M.A. Quevedo-Lopez, H. Stiegler, B.E. Gnade, H.N. Alshareef, Org. Ele. 11 (2010) 925–932. [176] G.G. Lee, E. Tokumitsu, S.M. Yoon, Y. Fujisaki, J.W. Yoon, H. Ishiwara, Appl. Phys. Lett. 99 (2011) 012901–012903. [177] R.H. Kim, H.J. Kim, I. Bae, S.K. Hwang, D.B. Velusamy, S.M. Cho, K. Takaishi, T. Muto, D. Hashizume, M. Uchiyama, Nat. Commun. 5 (2014) 3583–3594. [178] J. Liu, Z. Yin, X. Cao, F. Zhao, L. Wang, W. Huang, H. Zhang, Adv. Mater. 25 (2013) 233–238. [179] Y. Ji, B. Cho, S. Song, T.W. Kim, M. Choe, Y.H. Kahng, T. Lee, Adv. Mater. 22 (2010) 3071–3075. [180] S.R. Bakaul, C.R. Serrao, M. Lee, C.W. Yeung, A. Sarker, S.L. Hsu, A.K. Yadav, L. Dedon, L. You, A.I. Khan, J.D. Clarkson, C. Hu, R. Ramesh, S. Salahuddin, Nat. Commun. 7 (2016) 10547–10551. [181] M.T. Ghoneim, M.M. Hussain, Electronics 4 (2015) 424–479. [182] Y.S. Oh, S. Crane, H. Zheng, Y.H. Chu, R. Ramesh, K.H. Kim, Appl. Phys. Lett. 97 (2010) 052902. [183] Y. Li, Z. Wang, J. Yao, T. Yang, Z. Wang, J.M. Hu, C. Chen, R. Sun, Z. Tian, J. Li, L. Q. Chen, D. Viehland, Nat. Commun. 6 (2015) 6680. [184] G. Tian, F. Zhang, J. Yao, H. Fan, P. Li, Z. Li, X. Song, X. Zhang, M. Qin, M. Zeng, Z. Zhang, J. Yao, X. Gao, J. Liu, ACS Nano 10 (2016) 1025–1032. [185] C.W. Nan, M.I. Bichurin, S.X. Dong, D. Viehland, G. Srinivasan, J. Appl. Phys. 103 (2008) 031101. [186] Y. Wang, J. Li, D. Viehland, Mater. Today 17 (2014) 269–275. [187] Y.J. Chen, Y.H. Hsieh, S.C. Liao, Z. Hu, M.J. Huang, W.C. Kuo, Y.Y. Chin, T.M. Uen, J.Y. Juang, C.H. Lai, H.J. Lin, C.T. Chen, Y.H. Chu, Nanoscale 5 (2013) 4449– 4453.

42

Y. Bitla, Y.-H. Chu / FlatChem 3 (2017) 26–42

[188] S.C. Liao, P.Y. Tsai, H.J. Liu, J.C. Yang, S.J. Lin, C.H. Lai, Y.H. Chu, ACS Nano 5 (2011) 4118–4122. [189] S.H. Xie, F.Y. Ma, Y.M. Liu, J.Y. Li, Nanoscale 3 (2011) 3152–3158. [190] Y.G. Ma, W.N. Cheng, M. Ning, C.K. Ong, Appl. Phys. Lett. 90 (2007) 152911. [191] Y. Zhang, C. Deng, J. Ma, Y. Lin, C.W. Nan, Appl. Phys. Lett. 92 (2008) 062911. [192] S. Ryu, J.H. Park, H.M. Jang, Appl. Phys. Lett. 91 (2007) 142910. [193] H. Ryu, P. Murugavel, J.H. Lee, S.C. Chae, T.W. Noh, Y.S. Oh, H.J. Kim, K.H. Kim, J.H. Jang, M. Kim, C. Bae, J.G. Park, Appl. Phys. Lett. 89 (2006) 102907.

[194] L. Yan, M. Zhuo, Z. Wang, J. Yao, N. Haberkorn, S. Zhang, L. Civale, J. Li, D. Viehland, Q.X. Jia, Appl. Phys. Lett. 101 (2012) 012908. [195] Z. Wang, L. Yan, Y. Yang, J.F. Li, J. Das, A.L. Geiler, A. Yang, Y. Chen, V.G. Harris, D. Viehland, J. Appl. Phys. 109 (2011) 034102. [196] N.M. Triet, T.Q. Trung, N.T.D. Hien, S. Siddiqui, D. Kim, J.C. Lee, N.E. Lee, Nano Res. 10 (2015) 3421–3429.