Structural details of Ge-rich and
silver-doped chalcogenide glasses
for nanoionic nonvolatile memory
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Phys. Status Solidi A 207, No. 3, 621–626 (2010) / DOI 10.1002/pssa.200982902
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applications and materials science
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Maria Mitkova* , Yoshifumi Sakaguchi , Dmitri Tenne , Shekhar Kumar Bhagat , and Terry L. Alford
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Department of Electrical and Computer Engineering, Boise State University, Boise, Idaho 83725, USA
Department of Physics, Boise State University, 1910 University Dr., Boise, Idaho 83725-1570, USA
3
School of Materials, Arizona State University, Tempe, Arizona 85287-8706, USA
4
JAEA, 2-4 Shirane, Shirakata, Tokai-mura Naka-gun, Ibaraki 319-1195, Japan
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Received 20 August 2009, revised 9 December 2009, accepted 9 December 2009
Published online 25 February 2010
PACS 61.05.cp, 61.43.Bn, 61.43.Fs, 64.70.P, 78.30.Ly
* Corresponding
author: e-mail mariamitkova@boisestate.edu, Phone: þ1 208 426 1319, Fax: þ1 208 426 2470
We are reporting our results of Raman and X-ray diffraction
(XRD) studies on amorphous Ge46S54 thin films and the films
after silver photodiffusion. Based on the Raman scattering
studies, a structural model for amorphous Ge46S54 is suggested
including the formation of single Ge–S chains with a vibrational
mode at 410 cm1. The result of XRD measurement indicates
´
that there exists a medium-range order with about 6 Å even at
such Ge-rich composition. After the introduction of silver, the
medium-range order is lost and there was a change in the
diffraction curve indicative of the change in the local atomic
order. The experimental results are explained in terms of our
structural model, in connection with the application for fast
switching memory devices.
ß 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
1 Introduction The fast development of semiconductor technology and the requirements for obeying the Moore’s
law bring the conventional silicon based technology to
unprecedented scaling constraints. The traditional memory
technologies are rapidly approaching miniaturization limits
as the industry moves toward memory cells with 22-nm
lateral features projected by the International Technology
Roadmap of Semiconductors for 2016 [1].The reason is that
they are based on charge storage – and it becomes
increasingly difficult to reliably retain sufficient electrons
in these shrinking cells. Magnetic and ferroelectric randomaccess memories have also serious scaling problems. The
new solutions emerging for nonvolatile memory beyond
2013 are based on resistance change rather than charge
storage. Among others, the major candidates in this field,
which have enough maturity to be considered by the
industry, include (i) phase-change memory in chalcogenides
(Chs). This is the most mature technology in developing for
more than 40 years; (ii) programmable-metallization-cell
memory in solid electrolytes. This is a very fascinating
technology, developing on a fast line; and (iii) resistancechange memory in transition-metal oxides [2]. Its application
is known but theory in progress. The three technologies are in
different stages of development [3] and will soon become
prospective candidates for commercialization.
Our recent research interests have been related to studies
of Programmable Metallization Cell (PMC) memory devices
and particularly the materials science related to them. The
PMCm is a high-density nonvolatile solid-state memory,
with low operational voltage, small power consumption, and
extremely good storage cell scalability. In it the information
is stored by the growing or dissolving of a metal dendrite in a
solid electrolyte. An oxidizable anode and an inert electronsupplying cathode formed in contact with a solid electrolyte
create a device that exhibits a polarity dependent switching
property. The intrinsically high resistance of the device can
be switched to a low resistance state by growing a stable
silver electrodeposit from the cathode to the anode. A reverse
bias dissolves the electrodeposit, causing the device resistance
to increase. PMC devices are quite simple and inexpensive to
produce. Furthermore, this memory technology meets the
requirements of the new generation of portable computer
devices by operating at a relatively low voltage while
providing high storage density and a low manufacturing cost.
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M. Mitkova et al.: Structural details of Ge-rich and Ag-doped chalcogenide glasses
The number of materials which can be utilized between
the two electrodes (the solid electrolyte) is enormously high
since the main requirements towards them are simple: (i) a
reasonably large conductivity of the cations of interest, and
(ii) an amorphous or defective crystalline structure with
channels for ion transport. Examples include Cu2S [4], SiO2
[5], WO3 [6], etc., However, the Ch glasses are the best
candidate for this because: (i) their not too highly
coordinated network allows the formation of channels for
metal link growth; (ii) the ions of the most used metals – Ag
and Cu have very high mobility in them; (iii) they are a good
supplier of electrons for the electrochemical process to
occur; (iv) the introduction of Ag or Cu in them results in
formation of phase separated structure containing metal Chs
with mixed electron/ion conductivity which allows fast
switching through connection forming between the metal
containing islands; and (v) they offer one more degree of
freedom for metal introduction through photodiffusion.
The Ge containing Ch glasses are the preferred candidate
for industrial application due to their thermal stability which
can handle high-temperature processing requirements. In
them Ge is usually fourfold coordinated. Report on their
Mössbauer spectroscopy data [7] show that their structure is
based on Ge–Ch tetrahedra (A) phase, ethane-like structures
with Ge–Ge bond, and three Ch connected to each Ge atom
(B) phase and (C) phase corresponding to outrigger raft (OR)
with a layered structure. The hosting Ge–Ch glass films are
usually formed in the industry through sputtering and, since
the targets can be made only from stoichiometric materials or
such enriched in Ge, the films are very Ge rich, usually
containing over 33 at.% Ge. These Ge very rich materials are
not well studied.
In the present work based on our Raman and X-ray
studies, we present our vision about the structural organization of the Ge very rich glasses which could be obtained
only in thin films. Further, the hosting Ch structure formed
after introduction of Ag in the films is discussed.
2 Experimental The Ge–Ch films were prepared by
thermal evaporation of Ge40S60 material using semi Knudsen
cell evaporation source. The film thickness was 50 and
300 nm. The compositions of the films were measured with
an electron probe microanalyzer in the system of the
scanning electron microscope (LEO 1430VP) using energy
dispersive X-ray spectroscopy (EDS). The results showed
that the films’ composition was Ge46S54. Raman spectra
were recorded using a Raman spectroscopic system of
Horiba Jobin Yvon T64000, in backscattering geometry. The
441.6 nm laser line of the helium–cadmium continuous wave
laser (Kimmon Koha Co., Ltd. IK5752 I-G) at a power of
73 mW was used for collecting the scattered spectra. Our
previous studies [8] show that within of 10 min of laser
radiation the films do not change due to radiation and we
made sure that the measurement time was well below this
time constrain. The X-ray diffraction (XRD) studies were
performed in Panalytical X’pert X-ray Diffractometer. The
particular measurement conditions were as follows: the XRD
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was carried out under 18 glancing angle with Cu Ka emission
l ¼ 1.5425 Å and a 2u range from 5 to 1008 with 0.18 step
width and 5 s/step. The data points have been averaged for
three experimental runs.
3 Results Figure 1 shows the Raman spectra of
Ge-rich Ge–S glasses (Ge 33, 36, 40, and 46%). The three
Raman spectra (a)–(c), are from the results of Takebe et al.
[9], which were measured using bulk samples. The spectrum
at Ge 46% was obtained from thin films prepared in the
present study. There are three main regions where peaks are
observed: below 175, 200–300, and 300–450 cm1. At Ge
33%, there are intensive peaks in the 300–450 cm1 region.
When Ge concentration becomes 36%, a peak in the 200–
300 cm1 region appears.
At Ge composition larger than 36%, a peak appears in the
region below 175 cm1. These features are consistent with
the results reported by Lucovsky et al. [10] and by Kotsalas
and Raptis [11, 12]. Among the peaks, the peak at 340 cm1
is attributed to the symmetrical breathing mode of S atoms at
Ge(S1/2)4 tetrahedron [13, 14] as the assignment is widely
accepted by many researchers. The intensity of this peak
decreases with increasing Ge composition. This suggests that
the number of the tetrahedral units decreases with increasing
Ge composition. However, the peak still exists even at Ge
36% and more, where the tetrahedral unit is not supposed to
exist according to the results of Mössbauer spectroscopy [7].
So, in the case of Ge 36% and more, we attribute the peak at
340 cm1 to the one of the vibrational modes of the ethanelike Ge2(S1/2)6 units, whose Raman active frequencies are at
240, 340, and 376 cm1 [15]. The presence of the ethane-like
Figure 1 Curve fit of the Raman spectra of Ge–S glasses. The dots
show the experimental Raman spectra with the background subtracted. The sum of the intensities of the peaks is indicated by a solid
curve in each figure.
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Phys. Status Solidi A 207, No. 3 (2010)
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units indicates the occurrence of Ge–Ge bonds, in other
words, chemical disorder in the system. The peak at
370 cm1 is often referred to as the companion mode Ac1 . It
is attributed by some researchers [16, 17] to the stretching
motion of the OR accompanied by S–S bond. Other authors
regard this peak as the modes to the vibration of S atoms on
the edge-sharing double bonds [18, 19]. In addition, there is
the vibrational mode of the ethane-like units, which is at
376 cm1. Jackson et al. [20] suggested from first-principles
molecular-dynamics simulations that there are two peaks
near 370 cm1; a peak at 373 cm1 attributed to the mode of
the edge-sharing cluster, and a peak at 366 cm1 related to
the mode of the ethane-like cluster. According to the results
of Mössbauer spectroscopy [7], at Ge33S67, 70% of the phase
is composed by tetrahedral unit (A phase) and 30% of the
phase is composed by ethane-like units (B phase).
Figure 2 shows the results of the XRD measurements for
amorphous Ge46S54 film and the film after silver photodiffusion. The curves were obtained by making difference
between two diffraction data with different film thickness
(50 and 300 nm). In the diffraction curve for Ge46S54 film, the
´
first sharp diffraction peak (FSDP) is observed at 1.0 Å,
which is close to the one reported by Fueki at al. [21] for
Ge40S60 glass. This indicates that the medium-range order is
still preserved in the Ge46S54 film. After Ag photodiffusion,
the FSDP vanishes as shown in Fig. 2b, revealing that the
medium-range order has been collapsed. In addition, there
seems to be a little change in the positions of the second
´
´
(about 2.1 Å) and the third peaks (about 3.5 Å) after the Ag
photodiffusion. This would suggest that the first neighbor
distance and the coordination number change after the Ag
photodiffusion. It is easily imagined that the diffusion of Ag
affects the coordination in the glass networks a lot and the
result is consistent with the expectation.
Diffusion of Ag causes serious changes in the molecular
organization as well, as shown by the Raman spectra in
Fig. 3.
Formation of Ag related thiogermanate structural units,
as also reported earlier [22], affects the Raman modes in the
range of the tetrahedral units. Since reduction of sulfur
content in the Ge–S network is expected to lead to increased
formation of ethane-like units, we studied the time evolution
of Raman spectrum for Ag photodiffused Ge30Se70 films
under the laser illumination with the wavelength of 441.6 nm
as shown in Fig. 4. According to Lucovsky et al. [15] and
Jackson et al. [20], the peak at 180 cm1 is attributed to the
vibrational mode of the ethane-like structure, while the peak
at 200 cm1 is attributed to the breathing mode of tetrahedral
unit. A small peak at 175 cm1 in the spectrum at 0 min must
indicate the presence of the ethane-like units. However, it
vanishes with time. Instead, the peak at 200 cm1 becomes
larger, shifting its position to a smaller wave number. This
would suggest the reorganization of the structure under the
laser illumination in Ag photodiffused Ge–Se films.
Figure 2 Diffraction intensity for: (a) Ge46S54 film; (b) Ag photodoped Ge46S54 film, obtained from the difference between two
diffraction curves with different thickness.
Figure 4 Time evolution of the Raman spectrum for Ag photodiffused Ge30Se70 film under the laser illumination.
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Figure 3 Raman spectra of Ge44S56 film (a) and Ag photodiffused
Ag: Ge–S film (b).
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M. Mitkova et al.: Structural details of Ge-rich and Ag-doped chalcogenide glasses
4 Discussion First we will discuss the structure of the
Ge rich films based on our Raman spectroscopy data. At high
Ge composition (35% and more) we could not fit the spectra
in the range from 300 to 450 cm1 without considering a
peak at 410 cm1, which becomes larger with increasing Ge
composition. The peak at 410 cm1 is not originating from
the shift of the 430 cm1 peak, because both – the 410 and
430 cm1 peaks were required to fit the spectrum at Ge 36%.
One could suggest that the increase of the peak at 410 cm1 is
related to the increase of C phase, which exists at high Ge
composition and is supposed to consist of double layer
structure in crystalline (c-) GeS7. But there is no peak with
such a high frequency in the spectra of c-GeS. We would
expect from the results, that there is a structural unit, which
has a stronger bond than that in the double layer, and it results
in such a high frequency vibration, in C phase. To the best of
our knowledge, there is no evidence or suggestion of the
existence of such a structural unit. In order to get a better idea
about the structural transformation occurring at higher Ge
concentrations from the Raman spectra, we have evaluated
the peak components as shown in the bar charts in Fig. 5.
In the figure, the intensity indicates the product of the
height and the width of the Gaussian peak. In the peak at
340 cm1, there are two components; the vibrational mode of
the tetrahedral units, which belongs to the A phase, and the
vibrational mode of the ethane-like units, which belongs to
the B phase. In the peak at 370 cm1, there are also two
components: the vibrational mode of the edge-sharing
tetrahedral units or the bond-stretching mode of S–S dimers
at the edge of OR, which belongs to the A phase, and the
vibrational mode of the ethane-like units, which belong to the
B phase. The peak at 255 cm1 contains only the vibrational
mode of the ethane-like units. It is considered that the ratio of
the peak intensities related to the ethane-like units,
I(255 cm1): I(ET, 350 cm1): I(ET, 370 cm1) is always
the same even if the Ge concentration of the sample changes.
Figure 5 The intensities of the Raman peaks; A, tetrahedral units;
; B, ethane-like units; ; C, layer-like structure ; black bars,
unassigned peak.
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The ratio can be obtained from the spectral peaks containing
no contribution from the A phase.
In Fig. 5, we assumed that the ratio is determined by the
spectrum at Ge 46%. With this assumption, the content of the
B phase increases with increasing Ge composition. But there
are some questionable points in the result: (1) There is a fairly
large number of the tetrahedral units even at Ge 40%. (2) The
edge-sharing tetrahedral structure units or S–S dimers
remain in the spectra at Ge 36%, and the content does not
change much from Ge 36 to 40%. (3) The content of C phase
does not increase from Ge 40 to 46%.
For the peak at 410 cm1, the intensity increases with Ge
composition regardless of the assumption. It increases
independently from the peak at 220 cm1. Therefore, the
structural origin of 410 cm1 must be different from the
double-layer structure.
In order to find out the nature of the appearance of such a
new structural unit with a vibrational mode at 410 cm1, we
have performed a virtual structural modeling. Figure 6 shows
the structural development of virtually made crystalline
Ge–S. High temperature crystal phase of GeS2 consists of
Ge(S1/2)4 tetrahedral units. There are streams of Ge–S
chains. Between two Ge–S chains, there are edge-sharing
tetrahedra. It looks as if the edge-sharing tetrahedra connect
the two Ge–S chains on both sides. Starting from this
structure, we subtract S atoms to make a Ge-rich compound.
Here, we assumed that the Ge–S chain structure is preserved.
By subtracting S atoms, the number of the tetrahedral units
decreases. At the beginning of the subtraction, the absence of
edge-sharing tetrahedral units is remarkable. This is natural
because removing one S atom among four S atoms in a
tetrahedral unit makes a loss of one tetrahedral unit while
removing one S atom among six S atoms in the edge-sharing
tetrahedra makes a loss of one tetrahedral set. The trend
coincides with the result of Fig. 1, in which the edge-sharing
tetrahedral set decreases rapidly from Ge 33 to 36%. Losing
the edge-sharing tetrahedral set affects the structure in the
middle portion between the two Ge–S chain streams. By
subtracting more S atoms, the number of the tetrahedral units
decreases further. As long as the tetrahedral units cling to the
Ge–S chain at a stoichiometric composition of Ge 33%, the
bond angle of S–Ge–S is supposed to be fixed to that in a
tetrahedron, 109.478. However, when a large number of
Figure 6 Structural development of virtually made crystalline
Ge–S material.
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Phys. Status Solidi A 207, No. 3 (2010)
tetrahedral units is removed from the Ge–S chain, the bond
angle in the Ge–S chain can change. Finally, the Ge–S chain
completely loses the tetrahedral units. At the middle portion
between the two Ge–S chains, a new Ge–S chain must be
formed. We would expect that the Ge–S chain, from which
the tetrahedral units are removed, and the newly formed Ge–
S chain at the middle portion combine together and form the
layer structure, which c-GeS possesses. In it, due to the
availability of heteroatomic chains, the formation of a
coordination bond is possible. In this case the lone pair of
electrons from the chalcogen is used for its establishment.
Recent ab initio molecular-dynamics simulations by Van
Roon et al. [23] have shown that the bond angle Se–Ge–Se is
about 90 8 which coincides with the possibility of formation
of such bond.
Now we have to connect the formation of this bond with
the newly appearing vibrational frequency at 410 cm1. One
important factor to determine the frequency is force constant.
Bond length can be a good indicator to compare the
magnitude of the force constant (although not always). To
see the relationship between the frequency and the bond
length, we have plotted these values as shown in Fig. 7. The
shorter bond approximately means the larger force constant.
Therefore, the plot indicates that a larger force constant
yields in a higher frequency. This is consistent with what one
can expect for a simple case of diatomic molecules. For the
Ge–S chains, we suppose that the frequency of the bondstretching mode of Ge–S chain is 410 cm1. The Ge–S
distance from XRD measurements is 2.14 Å [21]. The datum
point for selenium was obtained from the stretching mode of
amorphous selenium [18] and the Se–Se bond length in
amorphous selenium [24]. As one can see in the Fig. 7, the
data points for Ge–S chains and Ge–Se chains lie on a line
between the data points of chalcogen chains. This result
seems to support the idea that the peaks at 410 and 175 cm1
indicate the bond-stretching modes of Ge–S and Ge–Se
single chains, respectively.
625
Figure 8 Channels for Ag diffusion formed through breaking of
coordination bonds.
Presumably, such Ge–S chains behave as fundamental
units and the correlation between neighboring Ge–S chains
might be preserved in amorphous Ge–S system over a wide
Ge composition range. We attribute this to the origin of
FSDP. Vanishing of the FSDP after silver photodiffusion
indicates a breakdown of the correlation between neighboring Ge–S chains. Considering the intercalation of silver ions
into the space between two Ge–S chains by silver photodiffusion, the breakdown of the correlation is easily understood as shown in Fig. 8. In the figure, we assume that the
light illumination breaks Ge–S bonds and the broken sulfur
atoms are negatively ionized. In the case of Ge-rich Ge–S
films, a coordination bond, which bridges the gap between
two neighboring Ge–S chains, is weak so that it must be
broken by the light illumination. The sulfur ions would create
a channel, through which positively charged silver ions
diffuse. This specific structure regarded in terms of PMC
devices performance implies that these devices will be very
fast due to the creation in their structure of a channel for Ag
diffusion.
5 Conclusions We have studied Ge46S56 thin films
which are a paradigm for understanding the structure of very
Ge-rich Ge–Ch glasses. Through the analysis of Raman
spectra, we proposed the presence of single Ge–S chains and
coordination bonds in the Ge-rich Ge–S films. A light
illumination would cause the breakage of the weak
coordination bonds. As a result, silver ions are supposed to
be easily inserted between two Ge–S chains. Such situation
would give an advantage in the performance of PMC
devices. It is worth examining the performance using Ge-rich
Ge–S films.
Figure 7 Correlation between the position of the Raman peak of the
bond-stretching mode of the chalcogenide chain and the bond length.
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Acknowledgements The authors thank Phoseon Technology
for providing the UV LED system. Y. S. acknowledges support from
IMI-NFG (NSF grant no. DMR-0409588). D. A. T. acknowledges
support from DOE EPSCOR grant no. DE-FG02-04ER46142.
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