11.2.1 Generic five-layer ‘battery-type’ device design

Figure 11.1 is a sketch of a standard electrochromic device (Granqvist, 1995) with five superimposed layers on a single transparent substrate or positioned between two such substrates. The optical functionality originates in the electrochromic film(s) which alter their optical absorption when ions are inserted or extracted from a centrally positioned electrolyte. This transport is easiest if the ions are small, in practice being protons (H⁺) or lithium ions (Li⁺). The electrolyte can be liquid, solid inorganic or comprised of a polymer. The ions are moved in an electrical field between two transparent electrical conductors, and the needed dc voltage is around 1–2 V. Except for very small devices, external metallic electrical contacts (‘bus bars’) must normally be put over at least part of the circumference of the device in order to achieve a reasonably fast and uniform colouring and bleaching.

There are three different kinds of layered materials in the device: The electrolyte is a pure ionic conductor and separates the two electrochromic films (or separates a single such film from an optically inactive ion storage film). The electrochromic films are mixed conductors of ions and electrons, whereas the transparent conductors conduct nothing but electrons. Optical absorption sets in when electrons are inserted into the electrochromic film(s) together with the ions from the electrolyte and are localized on metal ions. The valence of these ions is then changed, and when the ‘extra’ electrons interact with the incident light they can acquire enough energy to jump across a potential barrier to a neighbouring metal ion site. The absorption mechanism is conventionally referred to as ‘polaron absorption’ in physics and as ‘intervalency absorption’ in chemistry (a somewhat more detailed explanation is given in Section 11.2.3 below). This simplistic explanation of how the electrochromic devices work indicates that they can be viewed as thin-film batteries with a charging state that corresponds to the intensity of the optical absorption.

It is possible, already at this point, to introduce a number of interesting properties of electrochromic glazings which make them highly relevant to eco-efficient building technology (Granqvist, 2012):

The electrochromic technology is not an easy one – which explains why it has taken so long to mature – and several more or less non-standard technologies must be mastered (Granqvist, 2008). Six challenges stand out for practical electrochromic glazings as listed below:

11.2.2 Practical constructions of electrochromic glazings

The scientific and technical literature show many examples of electrochromic glazings of various sizes and has done so for decades, but few of these examples can be considered as products ready for the market or even prototypes. A number of those that are currently (2012) being delivered to customers on a very limited scale, or at least shown to customers, are introduced next (see also recent articles by  Baetens et al., 2010a, and Jelle et al., 2012). All of these products or prototypes rely on electrochromic tungsten oxide films for at least some of the coloration.

There are numerous alternative electrochromic device designs as well, both based on an oxide-based ‘battery’ approach as in Fig. 11.1 and others. Considering first the ‘battery’ type, one may note that there are non-oxide inorganic electrochromic materials, and ‘Prussian Blue’ (de Tacconi et al., 2003) was mentioned above. Furthermore, electrochromism is a common phenomenon in organic materials, and a vast literature exists on this subject (Monk et al., 2007). Their durability under irradiation is much less than for the oxides, but the coloration efficiency (the change in optical absorption per unit of charge exchange) can be much higher in organic materials than in oxides. Metal hydrides represent another option and can display variable reflectance and can operate in conjunction with electrochromic thin films in devices; constructions with films based on nickel-magnesium hydride have been investigated in depth (Tajima et al., 2010). Today’s devices of this type tend to suffer from limitations in longevity, modulation span and high-temperature stability (Tajima et al., 2011).

There are also several device designs that are distinctly different from the ‘battery’ type. One of these is the suspended particle device (often referred to as an ‘SPD’) which is rooted in the pioneering work on ‘light valves’ done by Land already in the 1930s (Marks, 1969). Essentially, a suspension of rod-like molecules, for example of herapathite (Kahr et al., 2009), are aligned under an ac voltage of ~ 100 V (i.e., some two orders larger than for the electrochromic device illustrated in Fig. 11.1) and are randomly oriented in the absence of this field; the optical transmittance through suitably confined suspensions is then changed. The modulation can be strong for luminous radiation, but not for infrared light, and the devices exhibit some haze (Vergaz et al., 2008). Alternatively, liquid crystals can be used in several different ways to create variable transmittance; the most common construction with regard to glazings uses polymer-dispersed liquid crystals (known as ‘PDLCs’) and suffers from significant haze (Cupelli et al., 2009; Gardiner et al., 2009). Another possibility is offered by some organic compounds, which can display optical absorption when a small current is drawn through them, and this phenomenon has been used very successfully for some two decades in ‘self-dimming’ rear view mirrors for automobiles. As a final possibility one may point at reversible electroplating, which in principle is able to give variable reflectance between very widely separated limits. This option has been investigated intensely (Ziegler, 1999; Laik et al., 2001) but has not yet led to practically useful glazings.

11.2.3 Electrochromic thin films

There are two types of electrochromic metal oxides, which are referred to as ‘cathodic’ (colouring under ion insertion) and ‘anodic’ (colouring under ion extraction); they are discussed in detail in Granqvist (1995). The standard electrochromic device, such as the one shown in Fig. 11.1, embodies two electrochromic thin films and it is clearly advantageous to combine one ‘cathodic’ oxide (e.g., based on W, Mo, or Nb) and another ‘anodic’ oxide (e.g., based on Ni or Ir). Shuttling ions between the two electrochromic films one way makes both of these films colour, whereas shuttling ions the other way makes them both bleach; this is sometimes referred to as a ‘rocking chair’ operation. Coincidentally there are ‘cathodic’ and ‘anodic’ oxides which can work in tandem and jointly yield electrochromic glazings with a rather neutral visual appearance that is appropriate for general applications in architecture.

What is the origin of the electrochromism for these oxides? An approximate answer can be given by arguments based on the crystalline structure, and a detailed examination of the electrochromic oxides reveals that they can be represented as (defect) perovskites, rutiles, and having layer/block structures. All of these structures can be described as comprising ‘ubiquitous’ MeO₆ octahedra (where Me denotes metal) connected by sharing common corners and/or common edges. Edge-sharing is related to some degree of deformation of the octahedra. Only one electrochromic oxide falls outside this description and exhibits properties with both ‘anodic’ and ‘cathodic’ traits: this is vanadium pentoxide (V₂O₅) which can be viewed as built from square pyramidal VO₅ units. The octahedral coordination is very important for the electronic properties of the electrochromic oxides and leads to a qualitative model for the optical properties for all of the oxides mentioned above, as elaborated elsewhere (Granqvist, 1993, 1995).

The detailed mechanisms for the optical absorption in electrochromic oxides are often poorly understood. Generally speaking, the absorption is associated with charge transfer, and polaron absorption captures at the essential features (Granqvist, 1995; Niklasson and Granqvist, 2007). The electrons inserted together with the ions are localized on metal ions and, in the specific case of tungsten oxide, change some of the W^6 + sites to W^5 +. Transfer of electrons between sites designated i and j, say, then can be represented schematically as W_i^5 + + W_j^6 + + photon → W_i^6 + + W_j⁵⁺. This mechanism operates only as long as transitions can take place from a state occupied by an ‘extra’ electron to one available to receive that electron, and if the ion and electron insertion is large enough this is no longer the case so that ‘site saturation’ (Denesuk and Uhlmann, 1996) becomes significant. Electron transfer then can occur also according to W^4 + ↔ W^5 + and W^4 + ↔ W^6 + (Berggren et al., 2007). However, these latter kinds of charge transfer do not dominate since highly reversible electrochemical reactions limit the permissible insertion levels to those where W^5 + ↔ W^6 + are prevalent.

Mixed electrochromic oxides can offer a number of advantages, and by having a large variety of sites available for charge transfer, it is possible to achieve an increasingly wavelength-independent absorption (i.e., a more neutral colour). Other advantages of mixed oxides are the possibility to widen the optical band gap in order to give a higher bleached-state transmittance in nickel-oxide-based (Avendaño et al., 2004) and iridium-oxide-based (Azens and Granqvist, 2002) films, and to ‘dilute’ expensive iridium oxide without major effects on its electrochromism (Backholm and Niklasson, 2008; Harada et al., 2011). Still another advantage is that the coloration efficiency can be increased by mixing suitable oxides, as shown in recent detailed work on mixed tungsten-nickel oxide films by Green et al. (2012).

The ubiquity of the MeO₆ octahedra is important not only for the optical properties but also for the possibilities to accomplish facile ion insertion and extraction in the electrochromic oxides. This is so because the atomic arrangements give spaces – tunnels in three dimensions – that are large enough to serve as conduits for small ions. Furthermore, transition metal oxides can display many different types of crystallinity depending on temperature and pressure, and sub-stoichiometry can yield so called Magnéli phases (denoted W_mO_3m–1, with m being 1, 2, . . . for the case of tungsten oxide) with a combination of corner and edge sharing for the octahedral units.

Easy ion transport does not happen solely as a consequence of structures comprising MeO₆ octahedra but can occur as a result of film porosity ensuing from limited atomic movements during the thin film deposition process. In fact there is a large number of thin film technologies based on atomistic and particulate deposition, bulk coating and surface modification, as discussed elsewhere in this book (see Chapter 8). It seems that all of these technologies can be adapted, with greater or lesser difficulty, to the preparation of porous structures that are adequate for electrochromic oxide films.

11.2.4 Transparent conducting thin films

The transparent electrodes may be the single most expensive part in an electrochromic glazing and therefore they deserve particular attention. There are several categories of transparent conductors, each with its specific pros and cons (Granqvist, 2007; Ginley et al., 2010; Smith and Granqvist, 2010). This section gives a bird’s eye view of such films based on oxides, metals and carbon.

Heavily doped wide band gap oxides are commonly used as transparent conductors in electrochromic devices. The most pertinent materials are In₂O₃:Sn (ITO), ZnO:Al (AZO), ZnO:Ga (GZO), ZnO:In (IZO) and SnO₂:F (FTO), and the doping level is typically a few per cent. All of the materials can give a resistivity as low as ~ 1 × 10^–4 Ωcm, a luminous absorptance of only a few per cent in a film with practically useful thickness (300 nm, say), and excellent durability. The optical properties are very well understood from basic theory, which means that detailed and trustworthy simulations of the optical properties can be made for devices (Hamberg and Granqvist, 1986; Jin et al., 1988; Stjerna et al., 1994). Films of ITO, AZO, GZO and IZO prepared by well-controlled reactive dc magnetron sputtering onto glass and PET typically have resistivities of ~ 2 × 10^–4 and ~ 4 × 10^–4 Ωcm, respectively. High-quality films of FTO are normally made by spray pyrolysis onto the hot glass emerging from the leer during float glass production. Depositions onto flexible substrates introduces risks for cracking and accompanying loss of electrical conductivity, but this effect is generally not a serious one unless the bending ratio is as small as a few centimeters (Cairns et al., 2000; Lan et al., 2010). All of the mentioned oxides are transparent for most of the spectrum characterizing solar radiation.

ITO films can be expensive as a result of the high price of indium (despite the fact that this element is not uncommon in the earth’s crust) and require careful process control for deposition; zinc oxide-based films can have similar optical and electrical properties but usually demand even more stringent process control; and good FTO apparently can only be made on very hot glass. Thus each of the transparent conducting oxide films has particular challenges and there is today no ‘best’ alternative for applications in electrochromic glazings. Health issues for the manufacturing of transparent conductors have been brought to attention recently, and it has been reported that the production of indium-containing oxides may lead to pulmonary disorders sometimes referred to as ‘indium lung’ (Taguchi and Chonan, 2006), which clearly can be an important concern for large-scale manufacturing.

Metal films can serve as excellent alternatives to the oxide-based transparent conductors. The coinage metals (Cu, Ag and Au) have conductivities that are some two orders of magnitude higher than for the best transparent conducting oxides so that comparative electrical properties can be achieved at about a hundredth of the film thickness; the luminous absorptance of the metal films can be of the order of 10%. The metal films are stretchable to a much larger degree than the oxide-based films (Graz et al., 2009).

The relevant metal film thicknesses are extremely small, which means that details of the thin film growth are important. Continued deposition onto a dielectric substrate such as glass or PET causes the deposited metal to go through a number of distinct growth stages: tiny metallic nuclei are formed initially; they grow and create increasingly irregular ‘islands’; these ‘islands’ interconnect and form a contiguous meandering network at a thickness corresponding to ‘large-scale coalescence’; the network then transforms into a ‘holey’ film; and finally a well defined metallic film can be formed (Smith et al., 1986; Lansåker et al., 2009). The most interesting films have thicknesses only slightly above that for ‘large-scale coalescence’, in practice around 10 nm (Hövel et al., 2010). Reflectance at the two interfaces of the metal film limits the luminous transmittance to ~ 50%, but the transmittance can be very significantly enhanced if the coinage metal films are positioned between high-refractive-index transparent layers that serve as antireflection coatings for the metal films.

There are several alternatives to the oxide-based and metal-based transparent conductors that are explored today (Hecht and Kaner 2011), and carbon-based transparent conductors may be of particular importance. Thus meshes of carbon nanotubes can combine high transmittance for luminous and solar radiation with good electrical conductivity (Hu et al., 2010a; Niu, 2011). Another alternative – which currently enjoys intense interest – is graphene, i.e., atomically thin layers of carbon atoms arranged in a honeycomb lattice (Geim and Novoselov, 2007; Eda and Chhowalla, 2010); these layers can be prepared via mechanical or chemical exfoliation of graphite into individual sheets as well as by chemical vapour deposition. Successful roll-to-roll production of graphene coatings was first demonstrated during 2010 (Bae et al., 2010, 2012).

Metal-based nanowire meshes is another possibility, and it has been shown that silver nanowires with diameters of ~ 100 nm and lengths of ~ 10 μm can be produced in large amounts by inexpensive reduction of liquid silver nitrate (Hu et al., 2010b, 2011). Suspensions of these nanowires can be deposited, and the conduction between adjacent wires can be improved by annealing. These coatings can have good electrical properties but suffer from some diffuse scattering which limits their applications in electrochromic glazings.

A final example may be poly(3,4-ethylenedioxytiophene), known as PEDOT, which cannot quite compete with regard to performance with most of the other options mentioned above, but which nevertheless is interesting since it can be prepared by printing at a very low cost (Elschner and Lövenich, 2011).

11.2.5 Flexible electrochromic foil

Tungsten oxide is the most extensively studied electrochromic material and was discussed in some detail above. It has cathodic coloration and needs to be combined with an appropriate anodic oxide to create an optimized device. Iridium oxide works very well in tandem with tungsten oxide but it is one of the rarest elements in the earth’s crust and only found in abundance in a few places. An alternative is hence needed for large-scale applications, and hydrous nickel oxide is such a material as discovered in the 1980s by Svensson and Granqvist (1986).

A number of studies of electrochromic devices based on tungsten oxide and nickel oxide have been reported in the literature; they are either rigid and based on depositions onto glass (Mathew et al., 1997; Subrahmanyam et al., 2007; Huang et al., 2011) or flexible and deposited onto PET foil (Azens et al., 2003b; Niklasson and Granqvist, 2007). Figure 11.2 illustrates a specific device design that was mentioned in Section 11.2.2 above: one PET foil is coated with ITO and tungsten oxide, another PET foil is coated with ITO and nickel-based oxide, and the two electrochromic films are joined by an ion-conducting polymer adhesive. Figure 11.3 reports time-dependent transmittance at a mid-luminous wavelength of 550 nm for ten consecutive colouring/bleaching cycles. Colouring proceeds slower than bleaching and the values shown do not indicate the darkest state that can be reached. The highly repeatable properties should be noted, and the cycling can be continued for tens of thousands of cycles without severe loss of performance.

The desired transmittance interval depends on the intended use of the electrochromic device. For architectural glazings one may emphasize a high bleached-state transmittance, and then low-cost antireflection coatings may be of interest (Jonsson and Roos, 2010; Jonsson et al., 2010). However, other applications may require that a very dark state can be reached for the sake of glare control. A simple way to obtain a very small transmittance in the dark state is to put two or more electrochromic foils on top of each other (Granqvist, 2008). As a first approximation, the transmittance of a doublefoil device is the square of the transmittance of a single-foil device. Figure 11.4 indicates that the transmittance at a mid-luminous wavelength can approach 1% in a double electrochromic foil, while the bleached-state transmittance is still appreciable.

It was emphasized above that low-cost manufacturing is the key to large-scale implementation of electrochromic glazings in buildings. Figure 11.5 shows some initial results on the transmittance modulation in an electrochromic foil-type device prepared by roll-to-roll deposition onto a ~ 1-km-long and 0.6-m-wide PET foil.

11.3.1 Vanadium dioxide-based thin films: three challenges

Vanadium dioxide is an interesting and complex material with at least seven different polymorphs among which rutile VO₂(R), monoclinic VO₂(M) (Morin, 1959), and triclinic VO₂(T) (Mitsuishi, 1967) phases are similar in structure, and there are also tetragonal VO₂(A) (Oka et al., 1990), monoclinic VO₂(B) (Théobald et al., 1976), paramontroseite VO₂ (Wu et al., 2008) and body centred cubic VO₂ (Wang et al., 2008). The interesting thermochromic properties ensue from VO₂(R) and VO₂(M) which exhibit a reversible structural transformation and associated metal-insulator transition at a ‘critical’ temperature τ_C in the neighbourhood of a comfort temperature; VO₂(M) is semiconducting and reasonably infrared transparent at a temperature τ, so that τ < τ_C, whereas VO₂(R) is metallic and infrared reflecting for τ > τ_C.

The pioneering work on the thermochromism of VO₂ by Morin (1959) was performed on bulk specimens. However, it was soon realized that reactively sputter deposited and reactively evaporated thin films could exhibit a similar metal–insulator transition. Subsequently it has been found that virtually any thin film technology is capable of providing thermochromic VO₂. The reversibility of the metal-insulator transition can be excellent in films (Ko and Ramanathan, 2008), whereas bulk samples tend to deteriorate upon repeated thermal cycling around τ_c. The possibilities to create energy-efficient fenestration by letting solar energy into a building when there is a heating demand and rejecting solar energy when there is a cooling demand were pointed out already in the 1980s (Greenberg, 1983; Jorgenson and Lee, 1986; Babulanam et al., 1987), and various aspects of this technology have been reviewed several times more recently (Granqvist, 2007; Parkin et al., 2008; Saeli et al., 2010a,b; Smith and Granqvist, 2010).

Figure 11.6 introduces the characteristic features of the thermochromism that can be seen in a thin VO₂ film. Spectral normal transmittance T(λ) and spectral near-normal reflectance R(λ) were recorded at 22 and 100 °C, i.e., at τ < τ_C and at τ > τ_C. The data were obtained for a 0.05-μm-thick film produced by reactive dc magnetron sputtering as reported elsewhere by Mlyuka et al. (2009a). Similar optical data – usually of spectral transmittance – have been reported many times in the scientific literature and are hence very well established (see the paper by Li et al. (2012) for references). It is clear from Fig. 11.6 that the short-wavelength optical properties are similar irrespective of the temperature, while the infrared reflectance for wavelengths beyond ~ 1 μm is higher for τ > τ_C than it is for τ < τ_C, thus giving proof for the metal–insulator transition. The infrared transmittance at λ > 1 μim shows an analogous change.

Figure 11.6 also shows the spectral sensitivity of the light-adapted human eye, denoted φ_lum and lying in the 0.4 < λ < 0.7 μm wavelength range (Wyszecki and Stiles, 2000) and the solar irradiance spectrum for air mass 1.5 (corresponding to the sun standing 37° above the horizon), denoted φ_sol and extending across the 0.3 < λ < 3 μm interval (ASTM, 2003). One observes in particular that φ_sol drops sharply towards long wavelengths. Wavelength-integrated luminous and solar transmittance values are now introduced by

[11.1]

This equation was used to compute data, shown in Fig. 11.7, on T_lum(τ,t) and T_sol(τ,t) where t is the thickness of the VO₂ film (Li et al., 2010). These computations used standard formulas for thin-film optics (Born and Wolf, 1999) together with optical constants n and k of VO₂ as evaluated before (Mlyuka et al., 2009b). The discussion below will also make use of the corresponding dielectric constant, denoted ε = ε₁ + iε₂, which is related to the optical constants via ε₁ = n² – k² and ε₂ = 2nk.

The findings in Fig. 11.7 permit an identification of two properties of thermochromic VO₂ films that severely limit their applicability to energy-efficient fenestration:

A third limitation, obviously, is that the change of the solar transmittance should take place in the vicinity of a comfort temperature of ~ 25 °C, whereas

These linutations indicate three challenges that must be met for practical thermochromic glazings, and VO₂ has to be modified in order to achieve ∆T_sol >> 10%, T_lum >> 40%, and τ_c = 25 °C.

A moderate boost of T_lum and/or T_sol can be obtained by straightforward optical design using high-refractive-index dielectric coatings, and data are available for a number of two-, three- and five-layer coatings with vanadium dioxide films between films of, for example, titanium dioxide and zirconium dioxide. The best properties were reached for TiO₂/VO₂/TiO₂/VO₂/TiO₂ (Mlyuka et al., 2009a,b). However, these improvements are not sufficient for making thermochromic glazings of general practical interest. Fluorination is an alternative route to improve T_lum for VO₂ films to some extent (Khan et al., 1988; Khan and Granqvist, 1989).

11.3.2 VO₂ nanoparticle composites with enhanced solar energy modulation and luminous transmittance

‘Nanothermochromics’ is a new concept that was introduced recently by the author of this chapter and his coworkers (Li et al., 2010). It deals with VO₂ nanoparticles dispersed in a dielectric host and implies that the composite can be represented as an ‘effective medium’ with properties intermediate between those of the nanoparticles and their matrix. The particles are small enough not to cause optical scattering. The ‘effective’ dielectric function ε^MG is (Smith and Granqvist, 2010; Granqvist and Hunderi, 1977, 1978):

[11.2]

where ε_m accounts for the matrix and f is the ‘filling factor’, i.e., the volume fraction occupied by the particles. The calculations to be presented below employed f = 0.01 and a thickness of 5 μm (so that the VO₂ mass thickness was 0.05 μm, i.e., the same as for the VO₂ film reported on in Fig. 11.6).

Equation [11.2] is appropriate for the Maxwell-Garnett (MG) theory (Maxwell-Garnett, 1904), which pertains to nanoparticles in a continuous matrix (Niklasson et al., 1981). There are many effective medium formulations – applicable to a multitude of nanotopologies – all of which coincide in the limit of a small filling factor; this implies that Eq. [11.2] can be used in the present case without any loss of generality.

The parameter α in Eq. [11.2] is given by

[11.3]

where ε_p is the dielectric function of the particles and L is their depolarization factor. Spheres are characterized by L = ¹/₃. A random distribution of ellipsoids can be represented by a sum over three components – corresponding to the symmetry axes – to yield α in the dilute limit (Niklasson and Granqvist, 1984). The calculations presented below pertain to prolate (‘cigar-shaped’) as well as oblate (‘pancake-shaped’) spheroids with depolarization factors obeying ΣL_i = 1. The depolarization factors can be expressed in terms of the major (a) and minor (c) axes of the spheroidal particles through known formulas (Landau et al., 1984).

Wavelength-integrated transmittance values, specifically T_lum(τ,t) and T_sol(τ,t), were computed according to Eq. [11.1]. Data on ε_p were obtained from the literature (Mlyuka et al., 2009b), and ε_m was set to 2.25 which is applicable to a matrix of a typical glass or polymer. Furthermore, the spheroids were taken to be oriented at random. The free electrons in the high-temperature phase of VO₂ have an exceedingly short mean free path (Allen et al., 1993; Okazaki et al., 2006; Gentle et al., 2007) and, perhaps surprisingly, this innocent-looking oxide is not well understood in terms of fundamental theories. Notwithstanding this situation, the very small value of the mean free path makes it unnecessary to incorporate any size dependence on ε_p (as there is, for example, in coinage-metal nanoparticles (Granqvist and Hunderi, 1977) and for the non-uniform metal films mentioned above in Section 11.2.4 (Norrman et al., 1978)).

The upper panel of Fig. 11.8 shows the experimental arrangement wherein light with photon energy ħω (with ħ being Planck’s constant divided by 2π and ω being angular frequency) is incident onto the nanoparticles; this panel also defines an ‘aspect ratio’ as a relationship between the axial lengths by m = a/c for prolate spheroids and m = c/a for oblate spheroids. The data in Fig. 11.8 can be directly compared with those for a VO₂ film with t = 0.05 μm in Fig. 11.7. The middle and lower panels of Fig. 11.8 prove that the nanoparticle composites have much higher values of ∆T_sol and T_lum than the films, which clearly shows the superior properties of the nanother-mochromic composites. Spherical particles yield the highest transmittance. Still better properties can be obtained with VO₂ hollow nanoshells (Li et al., 2011a). In terms of the underlying physics, the infrared optical absorption of the VO₂-based nanoparticles is governed by a plasmon resonance that is present in the metallic state at τ > τ_C but absent in the insulating state at τ < τ_C (Bai et al., 2009; Li et al., 2010).

Nanoparticles based on VO₂ can be produced in numerous ways as briefly reviewed in recent papers by Li et al. (2010, 2011a, 2012). More or less symmetrical particles can be prepared by wet chemical methods, molten salt synthesis, confined-space combustion, etc. There are also many methods to make nanorods (prolate spheroids with large aspect ratio) and nanosheets (oblate spheroids with small aspect ratio) by these techniques. A metastable form, generally referred to as VO₂(B), can be produced by chemical techniques and is of much contemporary interest for electrical batteries; this material can be transformed into thermochromic VO₂. Furthermore VO₂ can be made by oxidation of metallic vanadium and by reduction of V₂O₅. Recent data are available also for vanadium dioxide nanoparticles doped with tungsten (Ye et al., 2010) and Mo (Chen et al., 2010); the relevance of this doping will be explained in Section 11.3.4 below.

The nanoparticle composites can be useful in several ways and may be incorporated in polymer foils or laminates for practical low-cost thermochromic glazing as indicated in Fig. 11.9. Another possible application is in composites of VO₂ and an oxide such as ITO (see Section 11.2.4), where the latter component can impart a low thermal emittance (Li et al., 2011b).

11.3.3 Mg-doped VO₂ films with enhanced luminous transmittance

The luminous absorptance of VO₂ is undesirably high, and Fig. 11.6 showed that T(λ) dropped distinctly at λ < 0.6 μm for a film thickness as small as 0.05 μm. This property can be understood as an effect of band-to-band absorption and hence is inherent in the material. This problem with vanadium dioxide has been known for decades, and much effort has gone into finding ways to diminish its influence. The problem received an at least partial solution in recent work on sputter deposited films (Mlyuka et al., 2009c), where it was shown that Mg doping could produce band gap widening, i.e., an enhancement of T(λ) for λ < 0.6 μm. The upper panel of Fig. 11.10 shows spectral data for a VO₂-based film containing 7.2 at% Mg at τ < τ_c and τ > τ_C and compares these results with corresponding ones for undoped VO₂. The lower panel of Fig. 11.10 elaborates the quantitative role of the Mg doping on T_lum, which is seen to go monotonically from ~ 40% to more than 50% when the Mg content is increased from zero to 7.2 at%. Some concomitant decrease of the distinctness of the thermochromism for λ > 1 μm is apparent from the upper panel of Fig. 11.10. Mg doping has a secondary beneficial effect as well, as it influences τ_c favourably, as mentioned briefly below.

A more detailed description of the band gap shifts in Mg-doped VO₂ films can be put forward via an evaluation of n and k for 0.3 < λ < 2.5 μm and T < t_c. The absorption coefficient α was then derived from α = 4πk/λ. Figure 11.11 shows an evaluation of (αħω)^1/2 versus ħω (as is appropriate for indirect allowed band gaps; Wooten, 1972). Two band gaps appear: one shifting from ~ 1.6 to 2.3 eV for rising Mg doping and another lying at ~ 0.5 eV irrespective of the Mg doping. A preliminary interpretation of the two band gaps is feasible from the band structure in the proximity of the Fermi level for VO₂ at τ < τ_c (Goodenough, 1971; Abe et al., 1997) as elaborated elsewhere by Li et al. (2012).

11.3.4 Doped VO₂ films with thermochromic switching at room temperature

Doped VO₂, denoted M_xV_1–xO₂, can exhibit shifted values of their ‘critical’ temperature, and, at least for bulk specimens, M being W^6 +, Mo^6 +, Ta^5 +, Nb^5 + and Ru^4 + yields a lowered τ_c, while M being Ge^4 +, Al^3 + and Ga^3 + produces an increased τ_c (Goodenough, 1971). Thin films of doped VO₂ display the same tendencies, but τ_c can also be influenced by several external parameters such as thickness, mechanical deformation, non-stoichiometry, small crystallite size, etc.

Tungsten is known to be the most efficient doping element for decreasing τ_c, i.e., it has the largest fall-off rate R as the amount of doping is increased. This material has been discussed in detail in the literature. For bulk crystals of W_xV_1–xO₂ it has been reported that R = 27 ± 1 °C/at%W (Goodenough, 1971; Hörlin et al., 1972), while R = 21 °C/at%W was stated in some subsequent work (Reyes et al., 1976). It follows that as little as ~ 2 at% of tungsten is able to bring τ_c to a comfort temperature. Thin films of W_xV_1–xO₂ also have decreased values of τ_c, but the quantitative magnitudes of R have been found to lie between 7 and 26 °C/at%W, i.e., differing by a factor of almost four; detailed data have been given elsewhere (Li et al., 2012). Tungsten doping has only a very minor effect on the optical properties (Tazawa et al., 1998).

There are two basic reasons why the data on τ_c for thin films are scattered, one being that the electrical properties depend strongly on crystallinity, grain size, and the conditions at the two interfaces of the films; the other reason is the difficulty of evaluating unique ‘critical’ temperatures from graded and hysteretic transitions of the resistance. Two experimental parameters of great significance are the substrate temperature during deposition and the post-deposition annealing temperature. Empirical data show that either of these temperatures must exceed ~ 450 °C (Li et al., 2012).

The films of Mg_xV_1–xO₂, discussed in Section 11.3.3 above, also have lowered values of their ‘critical’ temperatures, and R ≈ 3 °C/at%Mg describes the data (Mlyuka et al., 2009c). This may be compared with measurements on bulk crystals of Mg_xV_1–xO_2–2xF_2x, which yielded R ≈ 6 °C/at%Mg (Akroune et al., 1985).

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