!"" # $" " %%& $%& ' ( &
# ) *+-,'. M. Mitkova, M. N. Kozicki
Center for Solid State Electronics Research, Arizona State University,
Tempe, AZ 85287-6206, USA
Introduction
The marriage of silver and chalcogenide glass has been widely explored.
The importance of this combination was first established when in 1967 the effect of
silver photodiffusion in chalcogenide glasses was discovered [1]. In the following
nearly 40 years, great contributions have been made to the field [2-4] and the
photodiffusion effect has been applied in the fabrication of optical elements using
relief images [5,6], in micro-photo-lithographic systems [7,8] and in direct imaging
by photoinduced silver surface deposition [9,10]. The low free energy of
crystallization of silver (48 kcal/mol.) was a further reason to consider the
introduction of silver in chalcogenide glasses used for phase change optical
recording [11,12]. This enabled one of the main requirements for good optical
recording - high phase transformation rate. One other aspect of silver’s influence in
chalcogenide glasses that is of interest is the effect on the electrical conductivity of
the glasses, which can be changed by several orders of magnitude when silver is
introduced [13-15].
In order to utilize all the above-mentioned processes in technological
applications, one must have a strong basic knowledge about glass formation,
structure, and the main physicochemical properties of the materials. This will give
information on which compositions are suitable for particular needs and how we
may optimize the material structure for both ease of formation and performance.
Apparently, this important interconnection between applied and basic science is the
reason for the large number of publications that exist in the area of basic and
applied research on these systems, most recent review of which has been
performed by Frumar and Wagner [16].
In this chapter we will focus on the structure and properties of fourfold
coordinated silver-containing chalcogenide glasses. We will review and comment
on the wealth of scientific research performed on such materials and present an
interesting technological application in integrated optics. Processing procedure and
formation of model devices will be described as well.
Interaction between silver and the chalcogen elements
There is no question regarding what happens when Ag interacts with
chalcogens as the phase diagrams of Ag with all three chalcogen elements are well
known and show the appearance of a stable stoichiometric composition - Ag2Ch
[17-19]. Some results have also been obtained for other Ag - rich phases but as
their stability is vague and they have not been absolutely confirmed by the
investigations of chalcogenide glasses containing Ag, we will not discuss them.
The introduction of Ag into chalcogenide glass networks presents a major
challenge in terms of understanding the structure of the glasses, as the local
chemistry is not entirely clear. In addition, the spectroscopic methods that are
mainly applied to depict the structure of the glasses (e.g. X-ray or neutron
spectroscopy) give atomic distance but very often it is difficult to distinguish
between the first and the second coordination spheres and so the results can be
ambiguous. Even application of direct methods for characterization, such as
Differential Scanning Calorimetry (DSC), can be misleading when the resolution
of the particular analysis system is insufficient.
When discussing the coordination of Ag we can follow the arguments of
Kastner [20] who was the first to describe the bonding of metal atoms in
chalcogenide glasses. However, we have to consider the fact that in addition to the
covalent bond that is expected to form with the chalcogens, Ag offers three empty
s-p orbitals and is surrounded by the lone-pair electrons of its chalcogen neighbors.
The latter offer the opportunity for the formation of up to three coordinate bonds. A
coordinate bond is similar to a covalent bond and has similar strength but the
bonding electrons are supplied by one bonding partner (the chalcogen atom) [21].
So the lowest energy-bonding configuration for Ag in chalcogenide glasses is an
overall neutral complex with Ag positively charged with the negative charge
located on neighboring chalcogen or chalcogens [22]. The coordination of Ag
thereby can vary but the most expected value would be three - with one covalent
bond and 2 coordinated bonds. The results of the structural investigations of
different Ag-containing chalcogenide glasses that we will discuss later reveal this
main coordination of Ag. The opportunity for four-fold coordination also exists as
there is one more free s-p orbital at the Ag atom but evidently this does not satisfy
the requirements of the lowest energy configuration and the electronegativity of the
whole complex in the presence of other cations in the system so that the probability
of this type of bonding is lower.
One other approach to the silver coordination issue is to consider the
crystalline compositions, which are known to have low and high temperature
polymorph forms. Since in quenching processes, due to kinetic factors there is a
high probability that the high temperature form will be preserved [23] we will
consider this in our further discussions. In fact, the structure of the liquid that
solidifies in the case of Ag2Se has been found to be very similar to that of the high
temperature form [24] and this is one more argument in favor of this assumption.
In this case, a cubic structure has to be regarded - Fig. 1 in which the chalcogenide
anions form the b.c.c. unit, with the Ag ions rapidly diffusing across distorted
tetrahedral interstitial sites and through the trigonal interstitial sites, as the number
of the states that could be occupied by them is larger than the number of the Ag
ions. This structure is considered in detail in the next section in the context of glass
formation in Ag chalcogenides.
Glass formation in silver-chalcogenide systems
Theoretical aspects
The theoretical aspects of glass formation of Ag2Se have been discussed in
details by Boolchand, et al. [25] based on the Constraint counting theory [26]. It
postulates that the mechanically effective connectivity of a network can be lowered
when constraints are intrinsically broken. Boolchand [25] extends his discussion to
AgI, which possesses an ionicity of 0.77 on the Philips scale [27]. This ionicity is
close to the threshold value of 0.785 when the energy of Ag atoms taking on a
coordination number 6 is equal to that of ions with coordination number 4. This
enables the rapid diffusion of Ag ions. The I-Ag-I and Ag-I-Ag bond angles flex,
leading to the general recognition that the bond-bending constraint about both
cations (Ag+) and anions (I- and Se2-) must be intrinsically broken [25]. Assuming
that in the b.c.c. structure Fig. 1 of αAg2Se, on stoichiometric grounds an average
of 4 Ag atoms and 2 Se atoms are associated with the unit cell, thus requiring 8
tetrahedral interstitial sites to be populated by Ag atoms on the cube facesas these
are shared by two adjacent cubes. This leads to a coordination number of 4 and 8
for Ag and Se respectively. However, the mechanically effective connectivity of
such a network can be different as the number of independent bond-stretching
constants nbs is given by r/2 where r is the coordination number of the atom, while
the number of bond-bending constraints nbb is given by 2r – 3. This gives the total
number of constraints per atom, nc , as (5/2)r –3 [25]. Enumeration of bondstretching constraints (r/2) for both Se2- and Ag+ cations then gives a total number
of constraints for the 6 atoms in the b.c.c. cell of
2[nbs(Se)] + 4[nbs (Ag)] = 2[4] + 4[2] = 16
(1)
which yields 6nc = 16 or nc = 2.66. The mechanically effective connectivity r of
such network is than related to nc as follows:
nc = 5/2rm – 3 or
rm (Ag2Se) = 2.26
(2)
So, based on theoretical considerations and experimental evidence,
Boolchand [25] has shown the possibility for Ag2Se to vitrify due to its effectively
low mechanical coordination even though it has a high coordination number of 4 in
its cubic form.
When considering glass formation in Ag-chalcogenide systems, we will
focus directly on the three-component systems - first with Ag-bichalcogenide
glasses as the glass forming regions for these systems give us an idea about the
glass formation in the binary systems as well.
Fig. 1. Crystalline structure of cubic Ag2 Se [25].
Fig. 2 shows one combined picture of the glass forming regions in the three
possible combinations of silver and bichalcogenide systems [28 - 30].
Fig. 2. Glass forming regions in the three-component Ag – Ch systems [28].
Glass formation in the Se-S-Ag system
The Se-S-Ag system has the largest glass forming ability among the Agbichalcogenide systems [29]. This is mostly due to the fact that there is a great
miscibility between S and Se and that the two elements build solid solutions. In a
liquid state they form an equilibrium mixture from linear co-polymer chains and
cyclic eight-member monomer rings [31]. There are even data showing the
presence of mixed monomer rings with composition Se3S5 [32]. The glass forming
region is situated in the Se-rich region and includes glasses containing up to
20 at.% Ag. The coordination number (<r>) of these glasses is about 2.26, and is
lower than the optimal number predicted by the Constraint counting theory [33,34]
and the glasses are therefore underconstrained.
There are no particular structural investigations of these materials but we
assume that they are formed from a mixture of chalcogenide chains, ring
fragments, and Ag2 Ch units. The phase separation in the system is shown for some
glasses by electron microscope investigations but is not considered in the results
concerning Tg. Actually, silver could be regarded as a bridging element between
the chains, as recent Raman spectra reveal that the introduction of silver and iodine
brings about to a shift of the Se chain stretching mode towards lower wave
numbers (251 cm-1 to 236 cm-1 ) [35]. Indeed, the mode at 236 cm-1 corresponds to
the trigonal type of chain formation that barely exists in glassy Se [36]. Apparently,
the new-formed structure with the additives we investigate is compact and is
becoming more covalent with stronger inter-chain interaction. Depending on the
amount of additives, the breathing modes of Ag-Se bonds are also present in the
spectra, suggesting the formation of two types of structural entities that, due to the
small sizes of the species or because they form an continuous backbone, do not
affect the properties of these glasses on the macro scale.
There is one other peculiarity in this system that also should be discussed the fact that no glass formation has been found in S-rich compositions. One
possible reason for this phenomenon is the great difference in the melting
temperatures of Ag (961.7 oC) and the chalcogenide matrix (which varies from 70
to 111 oC). For these circumstances, Ag solidifies at a very early stage in the
quenching process and precipitates out of the glass.
The presence of silver drastically influences the electrical conductivity,
which is about 6 orders of magnitude higher for samples containing 20 at.% silver
compared to those without silver [29]. The nature of the conductivity is not
discussed in the published research but we assume that there is a considerable ionic
contribution since Ag is very mobile because of the long Ag-Se distance (2.66Å).
Accordingly, the Coulomb energy in these glasses (E C ≈ 5.5 eV) would become
appreciably small.
Wahab [37] points to interesting details in the conductivity of glasses with
composition S0.8Se16Ag0.2 that are characterized as having a very high preexponential factor (4.1 × 108) for the temperature dependence of the resistance. The
high value of C indicates that conduction indeed takes place in the extended states,
which is due to a decrease in the localized state density and a consequent decrease
in carrier scattering by lattice imperfections. The electrical activation energy for
this glass is about one-third of the optical gap. As already mentioned, Ag can form
coordination bonds with the chalcogens that have corresponding (empty)
antibonding levels, which could give localized acceptor states in the gap.
Apparently, the low electrical activation energy is related to formation of
coordination bonds between Ag and the chalcogens and respective acceptor states
in the gap.
Glass formation in the Se-Te-Ag system
The glass forming region in the Se-Te-Ag system is located within narrow
boundaries close to 100 at.% Se and comprises concentrations of 38 at.% Te and 17
at.% Ag [30]. Comparing the glass-forming region with the previously discussed
system – Fig. 2 it must be pointed out that substitution of S with Te in the ternary
systems causes a decrease of the glass-forming region, mainly concerning the
content of Te compared to S. This is in agreement with the generally accepted idea
that strongly directed covalent bonds play an essential role in glass formation. Te is
usually considered to be a poor glass former. This seems to be determined from its
tendency to octahedral coordination similar to the distorted cubic lattice type NaCl.
In this, the ratio of the radii of the first and second coordination sphere, r2/r1, is
1.99, while in hexagonal Se this ratio r2/r1 is 1.433 [38]. The high coordination of
Te is preserved in the molten state too and so the main coordination number of
these glasses reaches 2.4 at lower concentrations of Te. The higher coordination of
these glasses discerns itself also in the higher T gs compared to the previous system,
although the Ag - Te and Te - Te bonds have lower bond-strength energies (195.8
and 257.6 kJ/mol respectively).
In this system, the presence of Ag again causes an increase of the electrical
conductivity by about 4 orders of magnitude [30].
Glass formation in the S-Te-Ag system
As pointed out in [28] and shown in Fig. 2, no glasses have been formed in
this system despite the presence of some amorphous phases. We assume that the
reason is an integration of the specific features of S and Te that have been
discussed above.
Glass formation and properties of glasses containing
Ge-chalcogenides and silver
Before starting a discussion on these glasses, we would like to make one
remark regarding the nomenclature of the glasses that we will consider. Usually in
chalcogenide systems the cations are presented as “x” in the compositional notation
and the chalcogen is the remainder, i.e. 1-x. When we introduce a third element
(Ag), then its quantity is shown as “y” and the rest is the chalcogenide backbone.
So a glass containing Ag and Ge-chalcogenides (Ge-Ch) could be presented as:
(GexCh1-x)1-yAg y. In this way it is easy to keep track of how the composition of
the glass is related to the composition of the Ge-Ch backbone, whose features are
well known, and see the influence of Ag on the glass formation and on the
properties of the glasses.
Ge-S-Ag system
Glass formation
Conclusions from the data from investigations of this system are rather
uncertain. First, data of Kawamoto et al. [39] show the glass-forming region being
concentrated at higher Ge concentrations Fig. 3. Later, Feltz et al. [40] investigated
the same system and show one region placed almost in the center of the triangular
diagram with corners corresponding to GeS, GeS2 and Ag2S. This glass-forming
region was not confirmed by other authors, so we will not consider this data.
Recently Kawaguchi et al.[41] showed that in addition to the glass forming region
found by Kawamoto, there is also one other glass-forming region in the Se-rich
part of the triangular phase diagram - Fig. 3.
Fig. 3 Glass forming region in the Ge-S-Ag system [41].
In Ge-containing systems, glass formation is restricted by the higher fourfold coordination of Ge. Obviously, there are also some other reasons for the
specific shape of the glass forming region in this system and particularly the
formation of the corridor on the tie line GeS2 – Ag where no glass formation is
found up to a Ag concentration around 20at.%. The reasons for the specific glass
forming region in this system could be related to the compositional variation of the
structure of the glasses [42] in which nanoscale phase separation occurs [43,44].
Also close to the region of the hosting system where no glass formation occurs, a
stiffness threshold exists [45] and an intermediate phase is formed [46]. All these
considerations are a significant challenge to a better and complete understanding of
this system.
Structure
All structural investigations confirm that Ge is four-fold coordinated. For
Ag coordination, different details emerge using different techniques. The silver Kedge EXAFS experiments [47 and other works of this team referred there] and
other X-ray related studies show that at low Ag2S concentration the structural
surrounding of Ag is analogous to that of α-Ag2S, while at high concentrations the
structure is rather similar to α-Ag8GeS6. This suggests that there are concentration
regions where the glasses are built up by homogeneous clusters and there are
regions where phase separation occurs. However this is not stated in [46] and
literature thereafter but phase separation for low Ag2S content (0<y<0.1) has been
documented by TEM characterization [48] Fig. 4. This comprises two types of
clusters based on the GeCh2 crystal structure and Ag2 S aggregates with a β-Ag2S
network structure. For the Ag2S-rich glasses it is suggested that the structure
becomes homogeneous due to better solubility of Ag2S and homogeneous
distribution of Ag+ in the glassy matrix with a Ge - Ag distance around 3.8 Å,
which seems to occupy various distorted sites [49,50]. EXAFS investigations of the
Ag rich phases suggest fourfold coordination of both cations (Ge and Ag) [50].
This is likely since it would lead to formation of materials with high coordination
but lower effective mechanical mean coordination as we have seen from the
consideration of the glass-formation in Ag2Chalcogenides. Results for materials
containing around 40% Ag atoms are also reported in the structural data from
isotopic substitution neutron scattering [52-54]. The authors suggest the presence
of some Ag-Ag correlations. Here also no indication of the occurrence of phase
separation is given and about the reason for the specific shape of the glass-forming
region. However, we note that these methods give information about the neighbor’s
coordination and often it is very difficult to distinguish the data for the first and
second coordination shells. As a result one could not expect that these methods
would give the correct picture of the mesoscopic structure of the glasses.
The first data regarding T g for these glasses have been submitted by
Kawamoto et al. [39]. It is difficult to discern the role of the backbone because of a
lack of systematic handling of the synthesized compositions. In the Ge-rich region,
the Tg obtained is somehow higher than the T g of the backbone (327 oC vs. 297 oC)
and this could be an indication of the formation of units with higher coordination.
In other words, we could think of these glasses as being formed from uniform
clusters containing Ge and Ag cations with Ag having coordination higher than 3.
For S-rich glasses, there are no data and this is one challenging and important field
of investigation yet to be explored.
Fig. 4. TEM of (GeS2 )0.9(Ag2S)0.1 ternary glass [48].
Electrical conductivity
The electrical conductivity in this system has been extensively
investigated. However, there are problems in the interpretation of the data as the
structure is not seriously considered in these studies and all conclusions regarding
electrical properties are mainly related to the amount of Ag introduced and not
connected to the backbone in which Ag is incorporated. Actually, the backbone in
which Ag is introduced may influence the structure of the formed glass
considerably and one can obtain either heterogeneous or homogeneous glasses. The
different forms would give rise to different electrical conductivity. Indeed, the
overall structure of the glasses makes them “tight” or “loose” ion conductors [55]
and thus we have to consider this. The temperature dependence of dc conductivity
of (Ag2S)0.5 – (GeS2 )0.5 glass shows that it is thermally activated with an activation
energy of 0.34eV [13,14] Fig. 5. By contrast, the temperature dependence of the
total measured ac conductivity, ( ), is considerably weaker than the dc value
when the ac component is dominant, and has an effective activation energy of ~ 0.1
eV [13,14]. Such a large difference in temperature dependence of the dc and ac
ionic conductivities can be related to ionic diffusion process in an inhomogeneous
system – a concept that is not accepted by Owens et al.[14] but we believe is
correct as we can draw some similarity with the analogous glasses from the Ge-SeAg system which would bring us to an inhomogeneous medium.
Fig. 5. Temperature dependence of the ionic conductivity for glassy
(Ag2S)0.5(GeS2)0.5 obtained from impedance spectroscopy measurements,
both in the dc limit ( extrapolated to =0) and at high frequency (
=10 kHz ( ), 100 kHz ( ), 1 MHz ( ), 10 MHz ( )). The temperature
dependence of the frequency exponent, s, of the ac conductivity is shown
in the inset. [14].
As mentioned a priori, the role of Ag on the conductivity of these glasses is
not discussed in the context of the hosting backbone. Though its role is very
important. As pointed out by Elliott [56], depending on the local structural
environment, the conducting ions could have different diffusion pathways and
activation energy. However there exist some data for (Ag2S)x(GeS2)1-x for which it
is reasonable to talk about the influence of Ag on the conductivity of the glasses at
limited concentration gradients, as in this case in which the backbone material
maintains almost the same structure. For these glasses, as shown by Ribes at al.
[57], the conductivity increases by a factor of 20 when the mobile ion content is
increased by only a factor of 2, with an associated decrease of the activation energy
from 0.36 to 0.32 eV– Fig. 6. More detailed studies of conductivity and 110Ag
tracer diffusion measurements for a Ag concentration gradient from 0.008 to 25
at.% [15,58, 59] reveal a nine orders of magnitude increase of the conductivity with
increase of mobile ion content. The change in the conductivity is accompanied by a
large decrease of the activation energy from 1 to 0.4 eV. This effect is explained by
the fact that at least three different transport regimes can be observed [57]: i) below
the percolation threshold at x~30 ppm Ag, the glasses are ionic insulators with a
very low ion transport number tAg+ 0.1 – 0.2; ii) just above the percolation
threshold for mobile ions the glasses become predominantly ion conductors with
tAg+ > 0.6 – 0.8 at x 0.008 at.% and increases rapidly to 1 with x (this transport
regime is attributed to the critical region just above the percolation threshold); iii)
far above the percolation threshold at x > 10 at %. In this silver concentration
domain the ionic transport is no longer dependent on percolation pathway. Rather,
it becomes network dependent with a strongly correlated motion of the Ag+ ions.
This is actually a modifier-controlled domain. An illustration of these three
transport regimes is given [60] on Fig. 7 a-c [58] As pointed out by Angell [61]
these are the conductivity pathways, which although being controversially
discussed for glasses containing Ag+, do really exist.
x
Fig. 6. Compositional dependence for the conductivity of
xAg2S-(1-x)GeS2 glasses [57].
Triangular voltammetry of (Ag2S)0.5 – (GeS2)0.5 glass [62] shows that the
electro active species at the anodic sweep are not particularly mobile. This can be
explained by the oxidation of sulfur in the glass to the 0 state. On the other hand,
on a cathode sweep, reduction of silver ions is detected but these can reoxidize
during the reverse sweep. However, a train of many sweeps generally results in a
break of the electrolyte in the electrode vicinity due to formation of silver dendrites
in asperities in the material which grow with sweep number and the resulting
mechanical stresses cause the glass to break [62].
Fig. 7. Three transport regimes in silver chalcogenide glasses: (i)
electronic insulator region below the percolation threshold at x~ 30ppm
Ag; (ii) critical percolation domain just above x ; (iii) modifiercontrolled domain far above the percolation threshold [60].
Some detailed ion dynamic studies for 0.5Ag2S-0.5GeS2 glass in extended
temperature (20 – 600 K) and frequency (10 Hz – 60 GHz) ranges have been
performed by Pradel et al., indicating the presence of three conductivity regimes
[63] – Fig. 8 a, b The first regime is observed in the low frequency-high
temperature region (i) where the conductivity does not depend upon frequency and
it is thermally activated with activation energy of 0.34 eV – it actually corresponds
to the dc conductivity of the particular glass. A departure from the Arrhenius rule
has been registered only at high temperature when the temperature becomes close
to Tg; (ii) At high temperatures-intermediate frequencies, a depressive region
(where s1 0.5 and A is thermally activated with an
exists with ( )=A
activation energy E1 (1-s)Ed.c.) - these data are in good agreement with previous
studies [62,64]; (iii) In the third regime ( )=B
. At T<100 K, this can be
accounted for by an asymmetric double well potential model. The third regime at
high frequency (GHz range) and high temperature (>150 K) clearly shows a super
linear frequency dependence. These results indicate that different mechanisms of
ionic conduction [65] characterize these glasses depending on the temperature and
frequency.
Fig. 8. Temperature dependence of conductivity for glass
0.5Ag S–0.5GeS in the whole frequency range [63].
One further development from the impedance study of these glasses [66]
shows that the nature of the ion dynamics in them can be explained by applying the
concept of mismatch and relaxation (CMR) [67]. As has been found for
0.5Ag2S.0.5GeS2 glass, the CMR model spectra (resulting from a rather general
approach which does not take into account any structural features) and nonvibrational conductivity spectra are in good agreement with each other since the
CMR spectra reproduce the increase of conductivity with frequency [66]. Fig
Finally, we would like to mention that conductivity measurements are a
good tool to characterize materials but one has to critically evaluate the data since
very often the different parameters can not be connected as simply as supposed
[68]. The best way to interpret the conductive characteristics of Ge-S-Ag glasses
that pose many contradictions is to correlate them to the structure of the glasses
[69].
Ge-Se-Ag system
Glass formation
The glass formation in this system has been investigated by Borisova et. al.
[70] and Mitkova et. al. [71] Fig. 9. The latter researchers have made corrections to
the borders of the glass-forming region although the principles of glass formation
have remained the same. As in the previous case, the glass-forming region of this
system consists of two regions separated by a corridor on the tie line connecting
composition GeSe2 – Ag where there is no glass formation up to a Ag
concentration of 20 at%.
Fig. 9. Glass-forming regions in the Ag–Ge–Se ternary system. The Serich and Ge-rich regions are denoted by I and II respectively [71].
Structure
The glass formation and the amount of Ag that can be introduced in these
systems are naturally closely related to the structure of the glasses and in particular
the hosting Ge-Se glass which undergoes a rigidity transition [72], which affects
the glass formation. Experiments to reveal the structure of these glasses have
included X-ray scattering [73-75], neutron diffraction with isotopic substitution
[50], neutron scattering [76-77], Raman scattering [78], Differential Scanning
Calorimetry (DSC) [79], electrical conductivity [39], dielectric spectroscopy [80]
and optical microscopy [80]. They have also been modeled using molecular
dynamics (MD) simulations [81] and Monte Carlo (MC) method [82]. In spite of
this large database, debate on basic aspects of glass structure continues. For
example, neutron structure factors [77], vibrational spectroscopy [78], DSC [79]
and MC [82] results on the ternary glass composition, x = y = ¼ , was found to be
consistent with a homogeneous network structure. Based on this, the proposal of a
specific 19- atom cluster of Ge3 Ag4Se3aSe3b(Se1/2 )6 stoichiometry anticipated in
ref. [78 and 82] as the element of medium range structure of a glass at x = y = ¼
deserves comment. In this cluster, Ge and Ag take on a coordination number (CN)
of 4 and 3 respectively, while Se takes on three environments with Sea possessing a
CN = 4, Seb a CN = 2, and 6 bridging (CN=2) Se atoms contribute 3 atoms to the
19 atom cluster. A simple count of the global connectivity of this cluster yields a
mean coordination number <r> = 3. Given the scaling of Tg with <r> in Ge-Se
glasses [83], the expected Tg of a glass having <r> = 3 would be greater than 400ºC
- this far exceeds the two Tgs of 230 ºC and 290 ºC observed for this glass
composition [71]. Given the measured Tgs, it is unlikely that such a highly
connected fragment could form the elements of medium range structure of the glass
composition of interest. Furthermore, justification of such homogeneous structural
model is put in question since MD analysis [81] of the neutron structure factor does
not support such a model. Besides, this pattern gives no clue to the fact why there
is no glass formation in the corridor starting at GeSe2 up to 20 at.% Ag in the
glasses.
A dual model for the structure of these glasses based on MDSC analysis,
Raman scattering investigations, and Mössbauer effect studies, has been proposed
in [71,83] and we will discuss this in terms of three different types of glasses – Serich with 0<x<0.33; Ge-rich glasses with x>0.33, and stoichiometric glasses with
x=0.33.
Fig. 10. Total heat flow in MDSC scans of Agy(Ge0.20Se0.80)1-y glasses with
increasing Ag content showing the evolution of bimodal Tg values. In
these scans the low-T endotherm is identified with the base glass (Tgb )
while the high-T endotherm is identified with the additive Ag2Se glass
phase (Tga )—see the text for details [83].
According to these studies [71, 83], Ag in the Se rich glasses plays the role
of network modifier and in the Ge rich part of the glass-forming region it is a
network former. This assumption was first confirmed by MDSC analysis of the
glasses as two Tg’s in the Se rich part of the system have been found - one at 230o
C that remains constant throughout the whole region of the glass compositions and
has been assigned to Ag2Se Fig. 10, and another T g that slowly increases with the
addition of Ag to the glasses because the Ge - Se backbone becomes progressively
Se deficient Fig. 11.
Fig. 11. Tg (x, y) variation in (GexSe(1-x))1-y Ag y ternary glasses [71].
Inset shows Tg = f <r> trend in Ge-Se glasses [45].
The foregoing thermal results can serve as the basis to elucidate the
molecular structure of the ternary Se-rich glasses starting first with a glass
composition at x = 0.20 and then those of compositions x = 0.25. These
consequences are then correlated with results of other types of measurements.
Macroscopic phase separation in Agy( Ge0.20Se0.80)1-y glasses
The MDSC results provide evidence of bimodal Tgs in virgin
Agy(Ge0.20Se0 .80)1-y glasses. In these experiments, the low-Tg endotherm in the
180 ºC <Tg < 200 ºC range is found to systematically up shift while the second T g
appearing at a fixed temperature (= 230 ºC) is found to display an increasingly
larger Cp step at Tg with increasing Ag content of the glasses. Both these features
are suggestive of a macroscopic phase separation [71] of these glasses that can be
described by the following stoichiometric relationship,
Agy( GexSe1-x)1-y = (3y/2)(Ag2/3Se1/3) + ( 1-3y/2) GetSe1-t
(3)
In equation (3) the first term on the right hand side designates an Ag2 Se
additive glass phase and the second term the remaining base glass phase. In
equation (3) t designates the stochiometry of the base glass and is given by
t = x(1-y) / ( 1-3y/2)
(4)
Thus, for example, starting with a base glass stoichiometry of t = 0.20 at
y = 0, addition of Ag results in t = 0.24 at y = 0.25. It is for this reason that the base
glass Tgb systematically up shifts as Ag concentration of the glasses increases.
From earlier studies [84] we know that the base glass T g equals 180 ºC at x = 0.20
and y = 0, and is lower than the Tg of the Ag solid electrolyte glass phase (230 ºC).
Independent confirmation to this picture of macroscopic phase separation was
given by alloying Ag2Se rather than elemental Ag in GeSe4 glass in a more recent
study [25]. In these studies, it was found that the Tg up-shift of the base glass is
absent, thus confirming the stoichiometry of the additive solid electrolyte phase.
Taken together these thermal results were among the first [71] to suggest that
glasses in the present ternary are intrinsically heterogeneous. The observation of a
separate Tg for the additive phase in these studies suggests that it is of macroscopic
dimensions (a few microns), and for that reason can be observed in optical
microscopy with relative ease.
Macroscopic phase separation in Agy( Ge0.25Se0.75)1-y glasses
Before discussing results on the ternary glasses, it may be useful to
remember that in binary GexSe1-x glasses, compositional trends in Tg(x) are well
documented [84] and reveal that as Ge content increases from x = 0.20 to x = 0.25,
Tg increases from 180 ºC to 230 ºC. This has the interesting consequence that in
the titled ternary, at low Ag concentrations (y < 0.10), it becomes difficult to
separate contributions to the glass transition endotherm of the base glass
(Ge0.25Se0.75) from those of the Ag2Se additive glass phase since their Tgs are the
same (230 ºC). The observed Tg of 230 oC for the Ag2Se glass phase suggests a
mean coordination <r> = 2.5 if one assumes close similarity of the dependence of
Tg on the coordination number for Ag2Se chalcogenides and the Ge-Se system.
This mean coordination number has been suggested for this material also by
EXAFS investigations [85] and is very close to the mechanically effective
coordination number of 2.26 discussed by Boolchand [25]. However as the
concentration of Ag additive increases further (y > 0.15), enough Se–deficiency of
the base glass sets in (equation 4) and drives up T g of the base glass so that it can
be discriminated against the Tg of the Ag2Se phase. The result is illustrated in Fig.
12a, wherein bimodal T g s are now observed in virgin glass samples at y = 0.20. At
still higher concentrations of Ag, such as y = 0.25, Tg of the base glass up shifts
enough to almost coincide with T x1, making it difficult to observe the second T gb
endotherm as in quenched samples. The controlled crystallization experiments
demonstrate that it is possible to observe glass transition of the base glass as shown
in Fig. 12b and 13b.
Fig, 12. MDSC scans of an Agy(Ge0.25 Se0.75)1-y glass at y = 0.20 in the (a)
virgin and (b) partially crystallized state. Upon heating at the indicated
crystallization temperature T1 x , the low- Tg a endotherm is suppressed
and the high-Tgb endotherm is enhanced. The sizes of the reversing heat
flow steps for the ‘a’ (additive) and ‘b’ (base) glass phases are related to
the contents of these phases in the glass of interest [83].
Fig. 13. MDSC scans of an Agy(Ge0.25Se0.75 )1-y glass at y = 0.25 in the (a)
virgin and (b) partially crystallized state. The Tg a (=230 .C) endotherm is
observed in the virgin sample, while the (Tgb = 290 .C) endotherm is
observed in the partially crystallized sample [83].
Because of its sheer simplicity, it is attractive to consider a model of glass
crystallization based on equation (3). Indeed, if equation (3) were to describe the
partial crystallization of the glasses upon heating to Tx1, one would expect the base
glass T gs to systematically up shift from 230 ºC at y = 0 , to 320 ºC at y = 0.25.
This is the case because equation (4) fixes t and the corresponding Tgbs can be read
in Fig. 2 of ref. [84]. In Fig. 14, the predicted increase of Tgb based on such a
model is shown as the broken-line. Although the predicted Tgb behavior parallels
the observed non-linear variation, it nevertheless systematically overestimates it,
suggesting that the model is not complete.
Fig. 14. Evolution of bimodal glass transition temperatures (Tga , Tgb )
upon increasing the Ag content in Agy(Ge0.25Se0.75)1-y glasses. The broken
curve gives the anticipated behaviour of Tgb if equation (4) is used to
describe phase separation of the ternary glasses. The dot–dashed curve
gives the Tg b (y) dependence as predicted by equation (6) [83].
As noted earlier, x-ray diffraction results reveal that the crystalline phase
that nucleates at Tx1 is c- Ag8GeSe6 and not c-Ag2Se [79,86]. To understand the
partial crystallization of the glass, it would be more appropriate to require the
following stoichiometric relation analogous to equation (3) to apply, i.e.,
Agy( GexSe1-x)1-y = (15y/8)( Ag8/15Ge1/15Se6/15) + Get’Se1-t’
(5)
where the first term designates the crystalline phase nucleated at Tx1 and the second
term is the remaining base glass. The Ge stoichiometry (t’) of the remaining base
glass is now given as
t’ = [x(1-y) –y/8]/[1-y-7y/8].
(6)
Equations (4 and 6) unequivocally show that for a fixed value of Ag alloying y,
t’< t, i.e. the base glass Ge stoichiometry for equation (6) is somewhat smaller than
for equation (4). Specifically, for a glass composition at x = y = ¼ for example,
equation (4) gives t’ = 0.29 while equation (4) gives t = 0.30. The predicted T g of a
glass at t = 0.30 is 325 ºC while that at t’ = 0.29 it equals 305ºC that is much closer
to the observed value (290 ºC). This is expected because crystallization of
Ag8GeSe6 depletes some GeSe2 from the base glass driving it Se-rich and thus
lowering T g. The predicted increase of Tgb (y) within such a model is thus in good
accord with the observed T gb (x = ¼,y) variation. The results from these ternary
glasses can therefore be understood well in terms of two separate glass phases and
the crystallization of c-Ag8GeSe6 at Tx1 Fig. 14.
The evolution of the Raman spectra of Fig. 15. is also very indicative of
the proposed structure [71]. Since Ag2Se is Raman silent; one can infer its presence
by the depletion of the Sen chain mode (CM) at 250 cm-1. At the highest Ag
concentration there is little or no scattering in Sen CM as all of the excess Se,
present in Se chains, is depleted to form the glassy Ag2Se phase that separates
from the GetSe1-t backbone (equation (3)).
Fig. 15. Raman scattering in (Ge0.25Se0.75)1-yAgy ternary glasses excited
by 647.1 nm line showing depletion of the Sen chain mode with Ag content
[71].
Ge-rich glasses
When we consider the Ge-rich Ge-Se-Ag glasses, we first have to bare in
mind the structure of the backbone. As reviled by the Mössbauer spectra, [87] in
the concentration region x = 0.32 to x = 0.43, 3 different types of structural units
develop in the structure of these glasses - Fig. 16. These nanophases are edgesharing Ge-Se tetrahedra (A), ethane like structural units (B) and units with
distorted rock salt structure (C) [87]. The last phase is the majority phase when
x > 0.40. In this phase the cations posses three long and three short Ge-Se bonds,
i.e., it is quasi three fold coordinated. In Fig. 17 (a) and 17 (b) the Mössbauer
spectra of a ternary (Ge0.4Se0.6)1-yAg y glass at y = 0.15 is compared to this with
y = 0. The observed line shapes reveal a partially resolved doublet. A
deconvolution of the resonance line shape in terms of two quadrupole doublets (B
and C) show that the site intensity ratio IB/IC increases from a value of 1.08 at
y = 0 to 1.44 at y = 0.15. Investigations at intermediate Ag concentrations confirm
the trend of IB/IC (y) increasing with y. These observations suggest that Ag
preferentially replaces Ge in the C- molecular phase, thereby precluding Sn
occupancy of it [83]. Ag also replaces Ge sites (quasitetrahedra) in the B molecular
phase, however such a replacement must lead to a severing of the Ge-Ag contacts
because of the Coulomb repulsion between the two electropositive cations. Such
chemical changes are thought to provide for a reduced connectivity of the
backbone (three long and three short bonds vs. three bonds) as reflected by the
decreasing of Tg [45].
Fig. 16. Mössbauer site intensity ratio, In/I(x), reflecting the concentration
variation of A, B, and C molecular phases in GexSe1-x glasses in the Ge
concentration range 0.32 < x < 0.40 [87].
Transmission, %
-7
-3.5
0
3.5
7
Velocity mm/s
Fig. 17. a, b Mössbauer spectra of indicated glass samples taken using
an emitter of 119mSn in CaSnO3 [71].
Stoichiometric glasses
In many respects, the glasses on the tie line GeSe2 – Ag, which we call
stoichiometric glasses, are the most interesting. We presume that the lack of glass
formation at low Ag concentrations is due to the lack of excess Se as all Se atoms
participate in Ge-Se tetrahedral structures and this causes Ag segregation in its
elemental crystalline form. In this region, bulk glasses form when y >0.20 as the
ternary melts become unstable against disproportion into Ag2Se and marginally
rigid Ge2Se3 nanophases, with Mössbauer spectroscopy providing evidence for the
Ge-rich phase. In the pristine GeS2 glass, two Ge(Sn) local environments appear Fig. 18 (a), and these were identified earlier [87] as Ge(Sn) in CS Ge(Se1/2)4
tetrahedra (narrow line -A and Ge(Sn) in ethane-like Ge2(Se1/2)6 units (doublet B). The qualitative changes in the line shape - Fig. 18 (c) observed upon alloying
Ag in GeSe2 once y>0.20, constitute direct evidence for the incipient disproportion
of melts (glasses) which was alluded to above. Specifically, we note that at x = 1/3
and y = ¼, equation (1) yields t = 2/5, i.e., the composition of the backbone is
Ge2Se3. It is for this reason that the spectrum of Ag alloyed GeSe2 glass in Fig. 18
(c) looks quite similar to that of the Ge2Se3 glass in Fig. 18 (b).
Transmission, %
Velocity mm/s
Fig. 18. Mossbauer spectra of GeSe2 with and without addition of Ag [71].
We suggest that the use of direct methods for investigation of the structure
of chalcogenide glasses such as spectroscopic and MDSC methods are most
informative for our understanding of the structure of these materials.
Neutron structure factors in Ag0.25( Ge0.25Se0.75)0.75 glass
The neutron structure factor Sn(q) of the titled glass composition measured
by Dejus et al. [78] has been analyzed by H. Iyetomi et al. [81] in MD simulations
using a 2-body inter-atomic potential and a system of N = 648 atoms confined to a
cubic box of 25.43 Å edge-length. The measured number density of the glass
composition fixes the volume of the chosen box. In the simulations, an MD
timestep of 0.004 psec was used, and melts thermalized at 1200 K for 20,000 time
steps, at 700 K for 40,000 time steps and at 300 K for 40,000 timesteps yielding a
total equilibration time of 0.4 nsec. In spite of the use of a short equilibration time,
and of 2-body forces only, the principal features of the observed structure factor
Sn(q) are reproduced by the simulations [81]. The pair distribution function T(r)
show that the principal Ge centered local units consist of Ge(Se1/2)4 tetrahedra in
the simulations, in accord with Raman scattering measurements [71]. The Ag-Ag
pair distribution function features [77] a broad peak centered around 3.5 Å and Ag
is thought to terminate the GeSe2 network to form Ag2Se-rich domains.
To decode the distribution of Ag in the simulations, the authors performed
a cluster-size distribution using a cut-off distance rc of 3.8 Å and found that Ag
atom distribution in the network is not random but actually clustered. The Ag
distribution result thus broadly supports a heterogeneous structure of the glass and
not a homogeneous one as was proposed by Dejus et al. [78] and Piarristeguy et al.
[82].
Electrical properties
After synthesizing the first Ge-Se-Ag glasses, Borisova et al. [70] provided
the first data regarding their electrical properties. They show a principal difference
in the variations of the electrical conductivity according to the amount of the
introduced Ag, which depends on the Ge-Se backbone in which Ag has been
introduced. While for the Se-rich glasses, there is a gradual tendency for
conductivity rise with the amount of introduced Ag and a change in the
conductivity type from electronic (p-type) to ionic has been found. When Ag is
introduced in Ge-rich glasses, their conductivity type remains electronic in the
entire concentration range of introduced Ag. Indeed, these results are in good
accord with the structure that we have obtained [71, 83] since in the case of Se-rich
glasses, due to the dominance of the Ge-Se backbone at low Ag-concentrations, the
glasses keep their electronic conductivity while at high Ag concentrations the
phase enriched in Ag dictates the conductivity. The lack of transition in the
conductivity type for glasses rich in Ge is apparently due to the fact that the
structure remains homogeneous in the whole concentration region and the related
electrical conductivity doesn’t change its character as Ag is chemically bonded to
the Ge-Se backbone. These results have been confirmed by later investigation of
the transport properties of the Ge-Se-Ag glasses [88] and dielectric loses [80]. The
impedance spectra measurements of Se-rich glasses Agx(GeSe3)1-x [89] show a
stepwise increase of the conductivity from 10-14 S/cm to 10-3 S/cm at x= 0.3 and the
understanding is that this rapid increase is chiefly due to appearance of silver ion
migration.
Formation of silver-doped Ge-Ch glasses by photodiffusion of
Ag in thin films
Silver photodiffusion – the fast diffusion of silver in a medium by
illumination with light, is a unique feature of chalcogenide glasses. Ever since this
effect was first encountered [1], it has been profoundly investigated as it can be
used in many applications of chalcogenide glasses. In the case of PMC
technologies, it is applied in the formation of the silver-doped chalcogenide glass
electrolyte.
Nature of the process
Our understanding is that the process of photodiffusion (here we have in
mind only the differences characterizing this compared to thermal diffusion) is
driven by the formation of charged defects in the chalcogenide glass, which form
on illumination with light and create an electrical potential. The light that is critical
for metal photodissolution is absorbed at or near the interface between the reacted
and unreacted (doped and undoped) chalcogenide layers [90, 91]. In this process,
holes are trapped by silver, while electrons move further into the chalcogenide film
and are trapped there. It has been shown [92, 93] that silver species move in the
doped chalcogenides as positively charged ions. The electric field formed by the
negatively charged chalcogen atoms and positively charged silver ions can be
sufficient for the silver ions to overcome the energy barrier at the interface.
Therefore the penetration of the metal into the chalcogenide during photodoping is
due to the difference in electrochemical potentials. The process of silver
photodiffusion was considered to be similar to that occurring in a galvanic cell,
where the more electropositive metal is dissolved into the electrolyte [94].
Kluge [95] considered the process of photodiffusion of metals in
chalcogenides as an intercalation reaction. The main reason this can be realized in
chalcogenide glasses is the fact that they possess relatively rigid covalent bonds
mixed with soft van der Waals interconnections. This type of structure ensures
formation of voids and channels where the diffusing ions can migrate and can be
hosted. The reaction can be efficient when the reversible transport of ions and
electrons can be achieved, accompanied by formation of bonds with the host
matrix, according to the reaction:
C20 + e- +M+
C1-M+
(7)
This reaction describes the transition of an initially twofold covalently
bonded chalcogenide atom (C20) into a C1- charged unit possessing only a single
covalent bond and an excess electron that establishes an ionic bond with Ag+ (M+).
Equation (7) shows the importance of the potential in forming the new C1-M+
bonds of another compound – the intercalation product. The possible number of
these bond-units is fairly high as the chalcogenide glasses are capable of forming a
number of single C1- centers under the influence of light illumination. Indeed, this
capability is a unique property of the chalcogenides and explains why this effect is
observed only in them.
Once silver is introduced into the chalcogenide glass, its further migration
into the chalcogenide glass continues. Lakshmikumar [96] suggested that the
doped–undoped region should be considered as being a semiconductor
heterojunction and proposed that it is the electric field at the p-n junction, which is
the driving force for the reaction. When the glass is illuminated with light, this field
is increased due to the formation of a number of defects. However there are
limitations to this consideration as this would be valid only in case when the
reaction product can be regarded as a doped semiconductor. Though, due to the
occurrence of high ionic conductivity, the interface should be considered as one
between a semiconductor and a solid electrolyte.
The photodiffusion kinetics depends on a number of factors such as light
intensity [85], light wavelength [97], temperature [98], pressure [99], external
electric field [100], composition of the hosting glass [101], and the atmosphere in
which the diffusion process is performed [102]. Many details in this respect are
given in the work of Kolobov and Elliott [3]. In this review, we will specifically
discuss the data concerning Ge-Ch systems.
Silver photodiffusion in the Ge-S system
The most profound investigation of silver diffusion in sulfur-rich Ge
glasses has been made by Oldale and Elliott [103]. They have found that the Ag
photo-dissolution rate in a-Ge29S71 has no induction period and the process has 2
stages – phase 1, which is an acceleratory stage leading to a maximum in the photo
dissolution rate, and phase 2, a final deceleratory stage. As the time to develop the
acceleratory stage shows a spectral dependency, it is obvious that the absorption of
actinic radiation in the photodoped film is responsible for this stage of the
photodissolution kinetic profile. Maruno and Ban [104] reported that when the
silver film is deposited on previously illuminated Ge30S70 film, the diffusion
process proceeds very slowly. Speculations are made upon the changes in the
structure of the chalcogenide film due to the light illumination. However, we
assume that oxidation of the chalcogenide film during the initial illumination could
also be responsible for the occurrence of this effect. The defects that can be created
by illumination are indeed the driving force for the oxidation process. Though they
can react with the atmosphere and form oxides as stated in [105, 106]. TEM
observations made by Maruno and Ban [104] indicate a distinct difference in the
microstructure and phase separation in their diffused films.
Kawaguchi and Maruno [107] obtained very important data about the
compositional dependences on the initial photodoping rate and photodoping
kinetics – Fig. 19 (a) and (b) Their data suggest that the kinetics of the diffusion
process is highly influenced by the chemical reaction between sulfur and silver in
the case of sulfur-rich glasses. A relatively low amount of silver is adopted by
these glasses and this fact should be considered in the context of the above results.
While in glass formation of bulk materials this is related to the large difference of
the thermodynamic constants of the two materials – the Ge-S backbone and the
formed Ag2S, in the case of thin films we assume this is due to the formation of
relatively large Ag2S clusters because the inhomogeneity of the hosting glass
prevents further diffusion. In the case of Ge-rich glasses, the high initial
photodiffusion rate could be related to the formation of the number of chemically
active photoexcited localized centers [108] that are responsible for the
photodiffusion process. The photodiffusion kinetics is also closely related to the
photoinduced changes in the hosting backbone [109] by which oxidation takes
place that affects the process as well.
When obtaining three-component glasses by silver photodiffusion in thin
films, one can introduce more silver into the particular Ge-Ch backbone than into
the bulk material. This has already been reported [110], where up to 60-67 at.%
silver was introduced in Ge30S70 thin films. As discussed by Kawaguchi et al. they
relate this phenomenon to the break-up of the Ge-S network structure through
formation of Ag-S bonds. Indeed, the first of these processes is part of the
photoinduced effects in these glasses by which the presence of silver can result in
dramatic changes of the backbone structure.
a
b
Fig. 19. a) Photodoping kinetics of Ag/GexS100-x samples (21<x<45). The
curves numbered 1-8 correspond to the time courses of samples of x =
21, 24.5, 30, 34, 35, 38, 42, and 45, respectively. The parameter y(t)
stands for the amount of dissolved Ag (y thickness) after exposure for
time t. The time course of x = 45 is shown by a dashed curve (8), since it
probably has a relatively large experimental error. Inset shows the
method evaluating initial photodoping rate (IPR) for the two different
types of data. b) Composition dependencies of IPR and total amount of
photodoped Ag, y ( ), at saturation level for Ag/GexS100-x samples The
data of x = 45 in parentheses probably involvee a relatively large
experimental error [107].
Silver photodiffusion in the Ge-Se system
The Ge-Se system has been more widely investigated than the previous
one. The main difference is that the films crystallize at lower silver concentration
and so introduction of silver is somewhat restricted compared to the glasses from
the Ge-S system. However the normalized comparative curves for the initial
photodoping rate and total amount of photodoped silver show great similarity [107]
Fig. 20 a) and b).
Fig. 20. Comparison of composition dependencies of ( ) Ag/Ge-S and (x)
Ag/Ge-Se systems for (a) y ( ) and (b) initial photodoping rate (IPR).
The curves are normalized by putting the maximum value to unity.
Comparison of composition dependence of Ag diffusion in Ge-S and GeSe [107].
Most extensive data about Ag diffusion in Ge-Se glasses has been given by
Kluge et al. [111]. They have found that the diffusion kinetics and the total amount
of diffused Ag are very closely related to the composition of the hosting backbone
Fig. 21. While for a backbone containing 75 – 90 at. % Se there is almost no
induction period for the diffusion process, at lower Se concentration this period
grows with decreasing Se. This is somewhat correlated with the depth profile of the
Ag diffusion. At high Se concentration, for GeSe5.5 , a step-like profile is found
[112,113] by which a Ag-depleted layer lies over the Ag-enriched layer situated
just above the substrate. Leung et. al. [114] also confirm that surface diffusion is
much smaller than the bulk diffusion. Indeed, some irregularities and discontinuity
of the silver doped film with formations of islands have been found also by Rennie
et al. [115]. For more Ge-rich glasses - GeSe3, the process is slower than in the
previous case [113] and follows the classical distribution. The same is valid for the
case when Ag diffusion occurs in the stoichiometric composition GeSe2 [116].
Wagner et al. [117] established a significant difference in the diffusion profiles of
laterally diffused Ag in Ge20Se80 and Ge40Se60 glasses, with sharp edge diffusion
for the Se-rich glass and a classical diffusion profile for the Ge-rich glass. Their
interpretation of this effect is related to the existence of 2 glass-forming regions in
the Ge-Se-Ag system. While for the Ge-rich glasses the diffusion process goes
through compositions characteristic only for one of the sub regions, in the case of
Se-rich glasses the composition of the diffused product resembles those of the 2
regions. This requires some structural rearrangements that affect the diffusion
profile. Since we know what the reaction products are when Ag is introduced into
the Se rich Ge-Se backbone, we assume that the chemical reaction that occurs is a
major factor in the acceleration of the process.
Fig. 21. Total amount of photodoped silver in dependence on exposure
time (parameter: Se content is 100- x) [111].
Structural investigations on the photodiffused material reveal formation of
a heterogeneous structure after photodiffusion [118,119]. Chen and Tai [118]
report formation of bcc Ag2Se, Ag2 SO4 and small amounts of orthorhombic Ag2Se
when Ag is diffused into GeSe2 glass and bcc Ag2Se and traces of free Ag when
the diffusion process is conducted in Ge0.1Se0.9 glass. We assume that these
differences in the diffusion processes are closely related to the availability of space
and channels for the diffusion of Ag in the particular hosting glasses. Formation of
Ag2Se has been submitted also by Zembutsu [119]. Considering the existing
results, Kawaguchi et al. [41] proposed a schematic model for the evolution of the
structure of the chalcogenide glasses during Ag diffusion that depicted the
formation of the two phases.
Optical programmable metallization cell (PMC) devices based
on Ag-Ge-Ch solid electrolyte thin films
Basic concepts of PMC devices
In recent years, a considerable increase in the application of solid-state
devices to the control of light has occurred. Many different device configurations
and materials have been used in such structures. Two major device types are liquid
crystal cells in liquid crystal displays (LCDs) and microelectromechanical systems
(MEMS) as optical switches. Although these two device types differ in many ways,
they both achieve the same goal of controlling the transmission and/or reflection of
incident light waves. In LCDs, light is selectively blocked or transmitted by the
liquid crystal cells. LCD cell production grew enormously with applications in
cellular phones, computer monitors, and other video applications, small and large.
Decreased size, weight and power consumption are among the main advantages
that LCDs offer but their main disadvantages are the manufacturing process is
awkward and they cannot easily be integrated. MEMS devices are becoming
widely applied in image projection systems and optical switches and routers.
Demand for these elements has increased with the implementation of digital light
processor (DLP) television monitors and state-of-the-art fiber telecommunications
networks that transmit large amounts of information with high speed. The main
disadvantages of these elements is that they are extremely difficult to manufacture
and the MEMS micro-mirrors require high voltage to function.
The functions of both LCD and MEMS switches potentially can be
replaced by an elegant device that is easier and less expensive to manufacture
based on Programmable Metallization Cell (PMC) technology. It offers a
revolutionary way to realize small, low power, and low cost opto-electronic
elements. The functionality is based on the formation of silver electrodeposits
using a solid electrolyte medium consisting of a chalcogenide glass doped with
silver. Depending on the programming field, the silver deposit is formed or
dissolved in or on the solid electrolyte by electrochemical means. With the growth
of the electrodeposit, the transparency in the channel over which it is formed
decreases dramatically; the electrodeposit, being metallic in nature, will reflect and
block incident light. By reversing the polarity of voltage the electrodeposit will
dissolve returning the channel to a transparent state. The geometry of the
electrodeposit is controlled by the duration and amplitude of the voltage applied as
well as by the composition of the solid-state electrolyte.
The solid-state electrolyte material formed by Ag diffusion in Ge-Ch
glasses can be used in PMC elements. This electrolyte will allow the movement of
silver ions under the influence of an electric field. If electrodes are formed in
contact with a layer of solid electrolyte, an anode, which has oxidizable silver and
an electrochemically inert electron-supplying cathode, a metal ion current flows in
the electrolyte when sufficient bias is applied (typically in the order of a few
hundred mV) and as long as there is oxidizable metal available from the anode.
The electron current from the cathode reduces an equivalent number of ions as
injected from the anode and a metal-rich electrodeposit is thereby formed in (or on)
the electrolyte. The amount of metal deposited depends on the magnitude and
duration of the ion current, i.e., the total Faradaic charge. The electrodeposit is
electrically neutral and the deposition process is reversible by applying a reverse
bias so that the electrodeposited metal is now the oxidizable “anode”. The reverse
ion current flows until the previously electrodeposited material is oxidized and
deposited back on the electrode, which originally supplied the metal. This
electrically stimulated injection and removal of metal in the solid electrolyte at
very low voltage is the basis of PMC memory technology [120, 121] but can also
be applied in other optical applications. These will be briefly discussed in the final
part of this chapter.
Formation of the solid-state electrolyte
We can form a good solid electrolyte, which will conduct ions over a wide
range of operational temperatures by photodissolving silver from a metallic surface
source into a thin film of germanium chalcogenide glass until the resulting ternary
is chemically saturated with the metal [122]. We have performed such experiments
using Ag-chalcogenide glass bilayers deposited by thermal evaporation on silicon
substrates. Complete photodiffusion usually occurs for Ag films a few tens of nm
thick after approximately 10 min of room temperature illumination with 405 nm
light, which is above the optical gap of the chalcogenide, with an optical power
density around 5 mW/cm2. Acceptable candidates for the hosting glass are GexS1-x
and GexSe1-x, where x is typically <0.33. These are chosen among the family of
chalcogenide glasses because of their high temperature stability, non-toxicity, good
electrical characteristics, and high silver diffusion rates. Silver uptake ceases when
all chemically available S or Se has reacted to form a separate phase close in
composition to Ag2S or Ag2Se. The amount of silver required for saturation
depends on the composition of the starting glass [122] but the resulting ternary can
have silver incorporation in the tens of at. %. Rutherford Backscattering
Spectrometry (RBS). Fig. 22 shows the change in the amount of photodiffused
silver with time in Ge20S80. The RUMP simulation demonstrates that the newly
formed composition at saturation for a composition Ge20S80 contains 47.3 at.%
silver – Fig. 23. The element markers denote the energy and the corresponding
channels when a specific element resides on the surface. Note that the distortion in
the S and Ge front edges and in the Ag back edges corresponds to enhanced
diffusion when compared to the overlaid simulation. In both cases, photo diffusion
and thermal diffusion, inspection of the Si signal reveals that the surface layer
contains approximately 5% Si.
Am ount of diffused Ag (at.%)
Fig. 22. RBS spectra of Ge30Se70 film photodiffused with Ag after 5 min
illumination and after 10 min illumination (the dashed line) [122].
Photo diffusion
50
40
30
20
10
0
0
5
10
15
20
25
Tim e (min)
Fig. 23. Amount of diffused Ag in Ge20Se80 as a function of the illumination
time as determined by RBS analysis [123].
A similar effect occurs for Ge30 Se70 starting material, although the
saturation amount in this case is lower (just over 33 at.%), as expected.
Characterization of the solid-state electrolyte
Silver photodiffusion causes complete rearrangement in the structure of the
hosting material as revealed by the Raman investigations [123] shown in Fig. 24.
Prior to diffusion, the hosting material exhibits modes of Ge-Se corner-sharing
tetrahedra at 194 cm-1, low scattering from edge sharing tetrahedra at 221 cm-1, and
Se chains at 260 cm-1 (stretching mode) and 150 cm-1 (bending mode). However,
after the diffusion process is complete, we see at 180 cm-1 the mode of the ethanelike molecule Ge2Se6 and at 203 cm-1 the mode of the GeSe4 tetrahedron. If we
assume that the mean coordination of Ag with Se is 3 we can use equation (3)
proposed in reference [71] to estimate the composition of the glasses after Ag is
introduced. At x =0.20 and y = 0.47, we get t = 0.36. We assume that the process
reaches saturation because of the strict number of free and under-coordinated Se
Counts (a.u.)
atoms that are available for reaction with Ag. Also, Ag concentration is higher than
in bulk glasses because the large difference in the melting temperatures of Ag2Se
and the Ge-Se backbone allows Ag precipitation during the bulk glass quenching
process and this limits the amount of Ag in the Ge-Se backbone to 33at. % in this
case [71].
a)
Ge 20 Se 80
b) (Ge 20 Se 80 ) 53 Ag 47
100
200
300
400
-1
Ram an Shift (cm )
Fig. 24. Raman spectra of: a) Initial Ge30Se70 film; b) Film resulting after
saturation with Ag [122].
There is excellent agreement between our Raman data and the postdiffusion calculated backbone composition of 36 at.% Ge. As discussed by
Boolchand [84], in this composition the underlying molecular phase consists of
face-sharing quasi one-dimensional ethane-like Ge2(Se1/2)6 chain fragments whose
presence is manifested in the Raman spectra by the appearance of the mode at 180
cm-1 depicted in Fig. 24 (b) following silver diffusion. This structure supposes the
appearance of Ge-Ge bonds. So far, investigations of the photoinduced changes in
Ge-Se glasses have not revealed the formation of Ge-Ge bonds. The act of
illumination of the film results in the creation of electron-hole pairs and is
accompanied by the formation of charged metastable states at the chalcogen [124]
that can react spontaneously with the surrounding Ag ions. This reaction will be
preferred as the energy that it requires is about 3 times less than the energy for GeSe bonding (202.5 kJ/mol vs. 484.7 kJ/mol) and we suggest that this is the reason
that more Se is consumed for formation of Ag2Se than in the initially available free
Se chains. This rapid reaction could be also the reason for the lack of an induction
period in photodiffusion.
XRD investigation shows that prior to any treatment, the bilayers consist of
microcrystalline Ag with 2 theta peaks at 38.12, 44.30, 81.54 and 97.89 degrees
(JCPDS card #87-2871) and an underlying amorphous matrix of Ge-Se glass, as
shown in Fig. 25. Following photo diffusion to completion, features of
orthorhombic Ag2Se with peaks at 2 theta 26.97, 33. 54, 40.41 - JCPDS card # 060501, 28.83 - JCPDS card # 25-0766 and 43.43 degrees JCPDS card # 20-1063
emerge together with peak of cubic Ag2Se at 35.89 degrees - JCPDS card # 27-
0619. From the XRD data, one can also see that the Ge-Se backbone remains
amorphous after illumination with light. Considering the spectra from the XRD
analysis, the unexpected presence of the high-temperature form of Ag2 Se, could be
due to space restrictions as Ag2Se forms in an existing solid-state framework. The
orthorhombic material has a more loosely packed structure than the cubic form and
we suggest that the surrounding hosting glass restricts its expansion. Although a
number of floppy units are available in the initial structure of the host, with the first
acts of silver inclusion and formation of Ag2Se, its structure becomes stiffer and so
the internal space limitation acts in the same way as elevated pressure, stabilizing
some clusters from the high temperature form which has the closest packing.
Norm alized intensity(a.u.)
10
8
6
+
4
Initial film
* *
++ ^ * *
++
*
*
*
*
2
20
30
40
50
60
70
80
90
100
2 Theta (deg.)
Fig. 25. XRD pattern of photodiffsed diffused Ge20Se80 film a) virgin film
b) film after illumination for 10 min.: * denotes peaks of Ag, + denotes
peaks of orthorhombic Ag2Se and ^ are peaks of cubic Ag2Se.
The X-ray Photoelectron Spectroscopy (XPS) analysis yields more
information regarding what kind of reaction products form as a result of thermal
and photodiffusion processes in Ge-S glasses [125]. The observed spectral shifts of
the S 2p peak obtained after Ag doping are reported in Fig. 26. Concerning the S
2p peak fitting, the effect of the spin-orbit splitting associated with a 2p core level
that gives rise to a doublet has been taken into account so S 2p1/2 and S 2p3/2 were
considered into the fitting procedure. Given that the films are evaporated on a
semiconducting substrate, there is a tendency for charging of these films, hence the
C 1s peak at 284.6 eV was used to correct for charging effects and Au 4d peak at
336 eV was applied for calibration of the experimental results. Both sulfur peaks in
the thermal- and photo-diffused cases are at lower energies than they would be for
pure sulfur (164 eV), indicating that sulfur exists mainly in the oxidized S2- state.
However, in the case of photoinduced Ag diffusion, this shift is 0.18 eV larger and
this pertains to a Ag2 S composition while in the case of the thermally treated
sample the data are more typical of the formation of GeS2 (see Table 1).
Table 1. Properties of the solid electrolytes used in the study [131].
Film
composition
occurring
after
saturation with Ag
Amount of
diffused Ag
at saturation,
(at.%)
Max.
amount of
Ag
in
bulk,
(at.%)
Sheet
resistance
of
Ag
doped film,
( /cm2)
Ge10.5Se42.3Ag47.2
Ge18Se42Ag40
Ge22.1Se44.9Ag33
Ge28Se42Ag30
47.2
40
33
30
33
32
20-30
25
2.2x10 3
4.3x10 4
1.2x10 6
1.6x10 8
Intensity (a.u.)
Ge20Se80
Ge30Se70
Ge33Se67
Ge40Se60
film
Binding energy
for elem ental S
Initial
composition
b)
Photo diffusion
a)
Therm al diffusion
155
160
165
170
Binding energy (eV)
Fig. 26. a. XPS derived S 2p electron spectra from the Ge-S:Ag film:
a) for thermally induced diffusion, b) for photoinduced diffusion [125].
Photo diffusion
1.06 eV
Elem ental Ag
Therm al diffusion
Photo diffusion
Therm al diffusion
Elem ental A g
Intensity (a.u.)
The data concerning the shifts occurring in the Ag 3d levels are shown in
Fig. 27. As one can see, while for the thermally treated sample the 3d peak appears
to be shifted by only 0.75 eV higher than in the case of pure Ag (368 eV), the
illuminated samples exhibit an additional shift of 1.06 eV and this strongly
suggests considerable formation of Ag2S following photodiffusion, i.e. there is
good agreement between these data and the XRD results given above about the
Ge-S system.
1.06 eV
b)
a)
365
370
375
380
Binding energy (eV)
Fig. 27. XPS derived Ag 3d electron spectra from the Ge-S:Ag film: a) for
thermally induced diffusion; b) for photoinduced diffusion [125].
The Auger Electron Spectroscopy (AES)-derived distribution of diffused
Ag as a function of depth in the Ge-S film is shown in Fig. 28. AES is a surface
sensitive technique and small amounts of typical contamination elements such as
carbon, oxygen and nitrogen are easily detected. In the case shown, carbon and
oxygen were detected but they are not considered here since they are not related to
the intrinsic nature of the effects investigated and do not affect them. Some amount
of substrate Si is seen in the initially sputtered film suggesting that the films are not
continuous. The reason for this could be some discontinuity of the film due to
shrinkage in the course of diffusion caused by the chemical reaction that occurs.
The slopes of the Si signals suggest that some Si diffusion also occurs into the
investigated films. This is consistent with the RBS data above.
100
Photo diffusion
Si
Com position (% )
80
Ag
60
40
C
20
S
0
0
10
Ge
20
O
30
40
Sputtering tim e (m in)
Fig. 28. AES determination of the relative change of the atomic
concentration in the Ge22S78:Ag film profile on silicon substrate as a
function of sputter time a) for thermally induced diffusion; b) for
photoinduced diffusion [125].
Silicon diffusion has never been noted in the investigations of the Si/GeCh glass interface. However recent results on solid-state diffusion on the Si – Ge
interface [126] have shown that Si tends to diffuse in vacancies in Ge and is also
capable of self-diffusion. Since chalcogenide glasses may occur as a medium
containing a high numbers of vacancies, one can understand the tendency of Si to
diffuse into the Ge-S film. We will not consider the effects related to Si diffusion
and contamination in the films, since they are not part of the particular effects on
which the present discussion is focused.
The AES data reveal that illumination with light causes penetration of the
diffused Ag, reaching about 80% of the film’s thickness, which in this case is 50
nm - (Fig. 28). There is an abrupt change in the slope right at the transition edge
with the chalcogenide film suggesting that there is no induction period involved
into the process, which agrees with what we concluded from the RBS analysis of
the film’s composition. When the diffusion process is driven by light, it has a
complicated character as the light affects the chalcogenide glass, creating a number
of charged defects. Davidova et al. [127] have demonstrated that for illumination
with a wavelength of 514.5 nm, which is close to the wavelength used in this work,
the bonding tendency of the free sulfur bond results in predominantly cis- rather
than trans-conformations. The bonds that occur are not discussed but we assume
that although some closed configurations are formed, a number of charged defects
occur that form a chemical potential for the reaction between diffused silver and
sulfur. It is for this reason that this process occurs immediately without an
induction period and the gradually developing chemical reaction drives silver ions
deeply into the chalcogenide film. Analogous results are also reported by Wagner
et al. [128] for silver thermal-and photodiffusion in As30S70 glasses where the
authors have found that the photoinduced diffusion results in about 3 times deeper
Ag penetration into the chalcogenide film. Also, a new structure of the films is
established after diffusion, where formation of Ag-S bonds is involved. The larger
chemical shift of the Ag 3d binding energy that is seen in the XPS spectra indicates
a very high ionized condition of the metal that could be the reason for an active
chemical interaction between the silver and the sulfur. This is confirmed by the
appearance of the shift to lower binding energies for sulfur. We assume that since
the chemical potential is the main driving force for the diffusion process, this is the
reason that great amounts of silver are introduced into the chalcogenide film –
saturation occurs at 43 at.%. This is some 3 at.% above that calculated for bulk
material and is a direct result of the changes induced in the hosting film by
illumination. The product of this process is mainly Ag2S, which according to eqn.
(3) leaves the hosting backbone richer in Ge. This is only possible because of the
formation of defects due to illumination with light. We suggest that the oxidation
processes in the system occurring during photodiffusion are predominantly
associated with the charged defects related to the germanium atoms, as the Ge 3d
binding energy is shifted towards 32.5 eV. This effect actually has been
documented in many previous studies of photoinduced phenomena in the Ge-S
system (see for example [3]).
Formation of silver deposits on blanket films
The structure of the blanket film optical elements consists of a Si substrate
covered with a layer of Si3N4 deposited by chemical vapor deposition, with a
thickness of 1800 Å, then 300 Å film of chalcogenide glass is deposited followed
by 150 Å of silver, both by thermal evaporation. The sample is then illuminated
with subband gap light to photodissolve silver into the chalcogenide film. The
excess silver is etched away using a solution of diluted Fe(NO3)3. Electrode metals
(Ag for the anode and Ni for the cathode) are thermally evaporated and patterned
using a standard photolithographic procedure.
Silver introduced in the chalcogenide films significantly modifies the
transport properties of the material so that ion mobility is relatively large, in the
order of 10-4 to 10-3 cm2/V.s at room temperature, and the availability of mobile
silver throughout the electrolyte is high [57, 41]. Silver is also used as the
oxidizable metal in the anode of the PMC structures and is particularly good in this
respect due to its nobility and the ease of both oxidation and reduction [129], so
that it is readily injected into and removed from the electrolyte. The resistivities are
2 mΩ.cm and 70 mΩ.cm for the selenide and sulfide silver-rich phases respectively
and is estimated to be as high as 108 Ω.cm for the germanium-rich phases. For a
silver-saturated ternary, the glassy interstices between the silver-rich regions have
an average width of less than 1 nm but the material’s high resistivity makes the
electrolyte resistance relatively high, in the order of 102 Ω.cm for the selenide-
based electrolyte and 103 Ω.cm for the sulfide-based material with the same silver
content.
The electron-supplying cathode is chosen to be electrochemically inert. A
wide variety of materials for this electrode can be used, ranging from refractory
metals to conducting oxides, but the results shown in this section relate mainly to
nickel cathodes.
After the device elements are formed, when a bias is applied so that the Ag
electrode is the anode and Ni is the cathode, all mobile silver in the current path
from the anode to the cathode takes part in the current flow and moves in a
sequential or coordinated fashion. At the nanoscale, each silver-rich region in the
current path acts as a local supply of ions. For each excess ion that enters one of
these regions from the anode side, one will simultaneously leave on the cathode
side and move into the interstitial zone there. Once in this glassy material, the high
local electric field will cause the ions to move toward the adjacent downstream
silver-rich region. The electron current from the cathode will also flow into the
electrolyte and the supply of both ions and electrons in the interstitial zones results
in electrodeposition so the excess metal in the electrolyte is effectively “plated” in
these nanoscale regions. The conductive electrodeposits bridge the interstitial
regions and help supply electrons to regions further away from the cathode until
the bridging to the anode is complete.
By using different glass compositions, materials optimization for
electrodeposit surface coverage can be determined. The electrodeposits are usually
grown with an applied voltage sweep from 0V to 5V and a current compliance of
1mA. Electrodeposits on Se rich Ge30Se70 and Ge rich Ge40Se60 glasses are shown
in Fig. 29 a and b.
The electrodeposition is limited by one of two factors; the time it takes for
Ag+ ions to reach the reaction side, and the time it takes for electrons to build up
enough charge to allow electrodeposition. For the particular case shown, it was
established that while for the Ge30Se70 glass the charge required for
electrodeposition was low - .551 S and the ion current limited the
electrodeposition reaction, for the Ge40Se60 glass the time needed to supply enough
electrons for electrodeposition is 9.662 ms – much longer than the time needed to
bring Ag+ ions to reaction site, 237.6 s, therefore the electron current limited the
electrodeposition reaction in this case. In all cases, when bias with an opposite
polarity is applied, the silver ions move in the opposite direction and the deposit is
destroyed, i.e. the medium reverts back to its initial optical properties.
Fig. 29. Microscopic view of grown electrodeposit: a) virgin film before
electrodeposition; b)
Morphology of silver deposits
The morphology of the deposits that grow when the electric field is applied
to the structure with coplanar electrodes is shown in Fig. 30(a)–(d) [130, 131]. One
can observe substantial differences in the shape and growth kinetics of the
electrodeposits and it is clear that they depend on the composition of the hosting
glass.
Fig. 30. Atomic force microscope analysis (3D topographical scan and
2D and line scan) of Ag electrodeposit grown on Ag-saturated: a)
Ge20 Se80; b) Ge30Se70; c) Ge33Se67; Ge40Se60 [131].
The general nature of the morphology of these deposits corresponds
closely to those reported in the literature in other systems, particularly those
formed by diffusion controlled processes such as diffusion-limited aggregation,
first described by Forest and Witten [132]. However the nucleation points from
which the dendrites start to grow seem to be related to the presence of excess Ag
ions in the electrolyte surface as in these regions the free energy for formation of
the electrodeposit will be lowest. In the case of Ge20Se80 glasses, we have to bear in
mind that the glass structure is floppy [84] and the illumination with light can cause
considerable depolymerization of the Se chains [3]. As a result, a number of
randomly distributed charge defects can occur in which the photodiffused Ag
reacts to form substantial Ag2Se regions that later act as nuclei for the formation of
dendrites during the process of electrochemical deposition. This is, in our view,
one reason that we see a great number of randomly distributed Ag dendrites with
small dimensions, covering the surface of the films and barely forming an oriented
morphology in the direction of the applied field. The electrons that have to be
supplied in order for Ag reduction to occur originate from the cathode and can flow
via the growing dendrite, enabling deposition at any point on the electrodeposit for
which there is a local supply of ions in the electrolyte.
Ge30Se70 glasses have a more stressed rigid structure [46] and illumination
with light cannot cause great redistribution of the existing Se chains so the charged
defects that form are closely related to the chain structure of the material. As a
consequence, some orientation of the electrodeposit occurs along regions, which
contain these surviving chains – Fig. 30.b. The same situation also holds for the
stoichiometric glasses – Fig. 30.c. It should be noted that these glasses have a
completely different type of conduction [70] and their overall conductivity is much
lower (Table 1). In the case of these materials Ag replaces some Ge atoms in the
backbone structure. We assume that localized Ag-containing units also occur in our
thin films and can serve as nucleation centers that are reasonably well insulated
from the rest of the film because of the heterogeneous nature of the hosting Ge–Se
glass [131]. These isolated nucleation centers are responsible for the disjointed
nature of the electrodeposits as the growth rate is much higher in these regions.
This same process is especially well expressed for the Ge-rich glasses, shown in
Fig. 30(d). In this case the hosting glass is thought to have n-type conductivity and
one can expect that it also supplies electrons in addition to the cathode for the
reduction of the Ag ions and this greatly accelerates electrodeposit growth, hence
much taller deposits can be formed.
Optical device testbed
The results from blanket film devices showed that the Ge40Se60
composition appears to be best suited for optical device development because of
the thickness and the surface area coverage of the electrodeposits grown. In our
optical device testbed, the silver doped chalcogenide glass is patterned to form bars
with different lengths and widths to confine the electric field and electrodeposits
grown. These have electrodes on either side, which are connected to probe pads via
planar interconnections; each of the columns of the device array share a common
anode but have individual cathodes. Each bar in the array therefore comprises an
individually addressable optical switching element. The array is shown in Fig. 31.
Fig. 31. Design of backplane testbed (5x3mm).
Fig. 32 and 33 show some data from representative devices, one with a
clear pixel and one with an occluded pixel in which the electrodeposit was grown
with an applied voltage of 5V. As shown, different degrees of coverage on the
active surface can be achieved. Note that it is apparent from this example that
silver electrodeposition on a solid electrolyte film formed by the photodissolution
of silver into a germanium-chalcogenide glass can be used to control the reflection
or transmission of light through an optical element.
Fig. 32. Clear pixel.
Fig. 33. Occluded pixel. Electrodeposit grown with 8 voltage sweeps. The
area covered is 45%. The electrodeposit was grown with 260 nC of
charge and had a volume of .65 m3, assuming a 100 nm height obtained
from AFM data.
The effects and processes depicted for PMC optical elements can be used
in a number of alternative optical devices, such as many of those described in other
chapters in this book. The richness and diversity in which PMC technology can
find applications in integrated optics arises from the fact that the elements formed,
by the use of transparent electrodes, could have vertical as well horizontal as
orientation and this doubles the opportunities for implementation.
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