Structural and Material Changes in Thin Film ... - IEEE Xplore

IEEE TRANSACTIONS ON NUCLEAR SCIENCE, VOL. 61, NO. 6, DECEMBER 2014

2855

Structural and Material Changes in Thin Film Chalcogenide Glasses Under Ar-Ion Irradiation Tyler Nichol, Student Member, IEEE, Muhammad Rizwan Latif, Student Member, IEEE, Mahesh S. Ailavajhala, Student Member, IEEE, Dmitri A. Tenne, Yago Gonzalez-Velo, Member, IEEE, Hugh Barnaby, Senior Member, IEEE, Michael N. Kozicki, Member, IEEE, and Maria Mitkova, Senior Member, IEEE

Abstract—We present results on structural and compositional changes in chalcogenide glasses under ion irradiation as a function of fluence and ion energies. Energy dispersive X-Ray spectroscopy (EDS) data obtained in this paper shows that the interaction with ions results in the loss of Ge atoms in Se-rich films. The compositional changes affect the structure of the films, which was manifested in differences observed in the Raman spectra. Ion interaction with of the films at the studied energies does affect the surface properties. Simulation of the penetration depth of the ions using Transport of Ions in Matter (TRIM) ions with the software shows that the interaction of incident chalcogenide glass occurs within the top 5-nm film thickness, with an etch rate for 450-eV ion energy of approximately 5 nm/s. We suggest the application of this effect for the formation of Redox Conductive Bridge Memory (RCBM) device arrays for which electrical characteristics are presented and discussed. Index Terms—CBRAM, chalcogenide glasses, ion beam radiation, memristor array fabrication, memristors, PMC, radiation-induced effects, TRIM simulation.

I. INTRODUCTION

C

HALCOGENIDE glasses are a segment of the amorphous glass family, which have unique optical properties in the IR region of the electromagnetic spectrum [1], [2]. These materials are highly transparent in this region and useful for many commercial, military and space applications. In some of these environments, primarily in space, these glasses can be exposed to energetic ions [3]. This exposure can change the structure of the glasses and alter the unique properties of these materials. One example for the application of these IR transparent materials is the James Webb Space Telescope [4]. Furthermore, there are many other conditions where chalcogenide glasses can interact with ions, for example ion implantation [5]. Studies have shown that deuterium irradiation causes relaxation and ordering Manuscript received July 10, 2014; revised September 22, 2014; accepted October 31, 2014. Date of publication November 25, 2014; date of current version December 11, 2014. This work was supported in part by the Defense Threat Reduction Agency under Grant HDTRA1-11-1-0055. T. Nichol, M. R. Latif, M. S. Ailavajhala and M. Mitkova are with the Department of Electrical and Computer Engineering, Boise State University, Boise, ID 83725-2075, USA. D. A. Tenne is with the Department of Physics, Boise State University, Boise, ID 83725-1570, USA. Y. Gonzalez-Velo, H. Barnaby and M. N. Kozicki are with School of Electrical, Computer and Energy Engineering, Arizona State University Tempe, AZ 85287-5706. USA. Color versions of one or more of the figures in this paper are available online at http://ieeexplore.ieee.org. Digital Object Identifier 10.1109/TNS.2014.2367578

of the amorphous network of chalcogenide glasses [6]. Using deuterium ions, as well as or ions, surface structures can be created [7], which by the application of appropriate materials and masks can lead to device preparation. This illustrates the multidimensional scenarios for which it is important to understand the nature of the effects that could occur. Additional interest towards this study arises due to the unique nature of chalcogenide glasses, which are a part of the disordered polymer system and their resultant response to ion bombardment should be significantly different when compared to the interaction of ions with crystalline materials [8]. One could expect scission or expansion of the polymeric structure due to the diffusion and incorporation of bombarding ions into it, ascribed to the free volume within this type of structure. Consequently, this could cause effects similar to the interaction of the chalcogenide matrix with other ionizing radiation such as or electron beam [9]. Most importantly, ion irradiation can prove useful in replacing sophisticated microelectronic processes for device formation as demonstrated in [7], holding a potential for preparation of memory devices based on chalcogenide glasses. However, for the production of devices with particular performance specifications, an in depth understanding the nature of all ion-induced processes is necessary, since ion irradiation, especially with heavy ions, can cause substantial changes in the electrical and optical properties of the chalcogenide thin films [10]. The above-mentioned properties of chalcogenide glasses are a function of their unique electronic structure. Chalcogen elements possess within their outer electronic shell, and elecelectrons are deeper into the electronic structure trons. The and they do not participate in chemical reactions with other elements. Two of the electrons form a lone electron pair, and the other two electrons contribute to the formation of chemical bonds with the surrounding elements. This electronic structure and the availability of the lone pair electrons contribute to the fact that all chalcogen elements have relatively similar electronegativity ( ), which leads to the formation of well-defined directional covalent bonds and satisfy the 8-N octet rule [11]. The low connectivity of the chalcogen elements which can lead to the formation of a two dimensional structure. However, the addition of other elements which tend to have a higher coordination, for example As or Ge, forms the well-established three dimensional network of the glass [12]. Here, we have to point out that this type of bonding between hetero-atoms is preferred in the chalcogenide glass structure because it reduces the structure’s enthalpy. The energetic ions that are present in the ambient space have very high energies. However, due to the

0018-9499 © 2014 IEEE. Personal use is permitted, but republication/redistribution requires IEEE permission. See http://www.ieee.org/publications_standards/publications/rights/index.html for more information.

2856


Fig. 1. RCBM device functionality.

Fig. 2. SEM image of the RCBM array.

shielding and coatings that are inherently available, these energies can become significantly attenuated to lower energies. This paper presents the effects of low energy ions on thin film chalcogenide materials, which could result after high energy ions penetrate through shielding and other materials. In this paper, we present data about the effects of ion bombardment on thin films from the Ge-Se chalcogenide glass system. This system, depending on the glass composition, offers a large variety of structural unit organization, including Se-Se chains (Ch), Ge-Se tetrahedral in corner-sharing (CS), edge-sharing (ES), and ethane-like (ETH) as well as layered structure (LS). Films with different compositions were exposed to three different ion energies with different exposure times to achieve different irradiation fluences. We have observed that there is a close compositional dependence of the effects. The studied structural and compositional changes arising in the chalcogenide matrix, as well as the simulated ion interaction, were applied to understand the ion-induced changes in the material. As a result of these studies we suggest a method for formation of Redox Conductive Bridge Memory (RCBM) device arrays [13] which can be integrated into neuromorphic or reconfigurable logic integrated circuits. The performance of these devices relies on the formation, low resistance state (LRS), and dissolution, high resistance state (HRS), of a conductive bridge between two electrodes, one of which is electrochemically inert and the other is based on a metal with high mobility, usually Ag. As a result of the diffusion of between the two electrodes, ), Ag following the electrochemical process (reduction of atoms deposit on the inert electrode. These Ag atoms deposit on this electrode until a Ag bridge between these two electrodes is formed. Fig. 1. illustrates the different conditions of device operation, such as the LRS and HRS states, that remain after the power to the RCBM device is removed, which classifies these devices as nonvolatile memory.

melt. The elements were melted in a Lindberg/Blue M rocking furnace reaching a temperature of 950 and quenched at 60 above the melting temperature of each composition. Thin film stack systems of ( , 0.3, 0.4) source material, with the evaporated compositions fluctuating from this composition, were prepared on a silicon substrate. First, 200 nm of was thermally grown on a Si substrate, followed by 100 nm of sputtered tungsten (W) using an AJA ATC Orion 5 sputtering system, and lastly m of chalcogenide glass was thermally evaporated using a Cressmbar vacuum. ington 308R coating system at Stack samples from each composition were bombarded with ions, having an initial ion energy of 150 eV, 300 eV, particles cm s, and 450 eV with a flux of to achieve a fluence of , and particles cm at each energy level. Ion bombardment was performed with a Veeco ME 1001 ion mill at an angle normal to the surface of each sample using a 300-mA beam current and chamber pressure of 0.2 mTorr. A nickel mesh with holes of nm nm was placed on a small area of the samples before ion bombardment, which allowed ions to pass only through the opening in the mesh, forming vias for RCBM devices. Using the same mesh, the silver electrodes were selectively deposited at the base of the vias by thermal evaporation using the Cressington 308R coating system. The newly deposited silver functions as the active electrode and W layer as the inert electrode, with chalcogenide thin film between them functioning as the active switching layer, thus forming a RCBM device array. Judging from the values of the conductivity, , depending upon the materials’ composition [14], these glasses are insulators rather than semiconductors. In this manner, the vias become naturally insulated by the chalcogenide glass between them, which confines the devices in the array, Fig. 2. Following ion bombardment, each sample was analyzed to determine its composition, chemical bonding, and surface roughness changes in the top chalcogenide layer and this data was compared to the as-deposited films. Particularly, energy dispersive X-Ray spectroscopy (EDS) was used to measure compositional changes caused by the ion bombardment, and also to determine the presence of Ar ions within the chalcogenide matrix, using a Hitachi S-3400N scanning electron microscope (SEM) equipped with Oxford Instruments Energy+ X-ray detector. Sampling five different areas on the samples,

II. EXPERIMENTAL Bulk chalcogenide glasses ( , 0.3, 0.4) were prepared from high purity germanium and selenium using a melt-quench method. Germanium and selenium were weighed to the appropriate atomic weights for achieving the desired compositions and placed in a quartz ampoule. The quartz ampoule was then placed under vacuum and sealed in preparation for the

NICHOL et al.: STRUCTURAL AND MATERIAL CHANGES IN THIN FILM CHALCOGENIDE GLASSES UNDER AR-ION IRRADIATION

Fig. 3. Film compositional analysis: (a)

; (b)

Fig. 4. Raman analysis of the evolution of (a) Se-Se bonding in radiation fluences with the respective energy.

; (c)

2857

.

. (b) Ethane-like bonding in

revealed that the average composition of the evaporated films and . as follows: , Raman spectroscopy was performed on five different locations on each sample contributing to the average and standard deviation values presented in Fig. 4(a)–(c). This spectroscopy method is useful to study the changes in the structure of the chalcogenide material. It was performed using a Horiba Jobin Yvon T64000 Raman spectroscopic system in back scattering mode. For excitation, a parallel-polarized 441.6-nm He-Cd laser was focused onto a circular spot of mm diameter at a laser beam intensity of 60 mW at cryogenic temperature and vacuum of mTorr. Although the laser wavelength is within the absorption spectrum of the films, no illumination-induced effects were observed during and after several measurements, which at the same time was enough for collecting a good Raman signal. The absorption coefficient for cm for wavelength of thin films is 441.6 nm, and increases with increasing germanium content [15]. The inverse of the absorption coefficient can be used to determine the penetration depth of the excitation light, and subsequently the probing depth of the Raman signal, which is approximately half of the penetration depth. For these measurements it is determined the measured Raman spectra originated from the top 28 nm of the irradiated samples. Previous investigations of one of us [16] show that at the applied evaporation technique the composition and the film’s structure in depth do not change, so that it is expected that all differences found by

.(c) Ethane-like bonding in

at various

the Raman and EDS studies will originate from the interaction of the films with the ion beam. The surface morphology of Ge-Se layer was studied using an OTESPA probe with Veeco Dimensions 3100 Atom Force Microscopy (AFM) system equipped with a Nanoscope IV controller in tapping mode. Since the interaction of the chalcogenide films with ions results in energy loss, which limits the penetration depth of the ions, ion bombardment simulations were performed using Transport of Ions in Matter (TRIM) software, which provides an insight into the depth of the ionization due to an incident ion. The target material thickness was modeled using the same thicknesses as the experimental films. Density of the film was derived from the densities of chalcogenide glasses as measured by Ivanov et al. as g/cm for [17]. To achieve a higher statistical average, 10,000 atoms were used, which gives a large certainty regarding the depth of penetration and distance of maximum ionization. III. RESULTS The EDS results revealed that the Ge concentration in the films incurred a significant reduction with increasing ion beam exposure, while the film concentration in the and films remains fairly stable as shown in Fig. 3(a)–(c). The Raman spectrum affirms the disordered nature of the studied films. The spectrum for Ge-Se system contains some

2858

Fig. 5. AFM 3D image of all films bombarded with ion bombarded samples. (c) (b)


particles cm Ar Ions at different ion energies. (a) ion bombarded samples.

intrinsic structures that are composition dependent. In the selenium-rich spectra, the structure is manifested through four distinct peaks located at cm , cm , cm and cm corresponding to corner-shared structure (Ge-Se-Ge), edge-shared structure (Ge-Se-Ge), selenium-selenium (Se-Se) vibrational bonds and the asymmetrically stretching edge-shared structures, respectively [18], [19]. In the Ge-richer samples ( , and ), ethane-like bonding (Se-Ge-Ge-Se) structure is exhibited in the cm . The Raman spectra were fitted Raman spectra at with Gaussian peaks, examples of which are illustrated in the insets of graphs in Fig. 4(a)–(c). Comparison of the areas of each peak revealed a unique trend in the Se-Se bonding and the ethane-like bonding, which are also presented in Fig. 4(a)–(c). In the films, there are two trends - first a decrease in the Se-Se bonding followed by an increase with additional ion fluence. In and films, the ETH-like structures predominantly react and their reaction is very complex, as shown in Fig. 4(b)–(c), which will be discussed in the following section. After irradiation, the films were also characterized using AFM to study their surface roughness as illustrated in Fig. 5(a)–(c). The surface roughness after the lowest radiation fluence is similar to the unirradiated samples. The Ge-rich samples that have been exposed to the largest ion energy for the greatest fluence exhibit a large change in the surface roughness when compared to the as-deposited films. Selenium-rich samples incurred a reduction in the surface roughness while the surfaces for the other two samples were significantly rougher after exposure as demonstrated in Fig. 6. A Wyko NT1100 optical profiler was used to determine the etch rate for ions with energy of 450 eV by measuring the step profile after bombarding the material through a mask for 30 s. The etch rate is approximately 5 nm/s which equates to etching 300 nm, 600 nm, and 900 nm at 60-, 120-, and 180-s exposures,

ion bombarded samples.

Fig. 6. AFM surface roughness for Ge-Se samples for different ion energies particles cm Ar Ions. bombarded with

respectively. The simulation result revealed that the maximum ionization and interaction of the incident ions with the chalcogenide glass occurs within the top 5-nm film thickness. This effect could be explained by the packing density within the hosting chalcogenide glass, data for which is presented in Fig. 7. [20]. Data regarding the molar volume was taken from [21] and packing fraction data was calculated according to [22]. Fig. 2. shows an SEM image of the fabricated -device array with Ge-Se film isolating individual cells at the maximum ion energy of 450 eV, which offers the highest throughput for devices fabrication. Experimental measurements revealed that the average thickness of remaining material at the base of the via was around 100 nm, with a standard deviation of 20 nm. A demonstration of the validity of this technology for fabricating arrays of devices has been achieved through their current–voltage (IV) characteristics, presented in Fig. 8. The Se-rich devices revealed a significant variability in the device


Fig. 7. Molar volume and packing fraction of Ge-Se glasses [20].

threshold voltage, while increasing the germanium content in the active film compositions resulted in devices demonstrating a consistent decrease in this variability. The devices with the highest Ge-content resulted in the most uniform threshold voltage and IV sweeps, presented in the inset of Fig. 8(c). IV. DISCUSSION The EDS data suggests that there is a loss of Ge-content from the films at 150 eV, which intensifies with the increase of the ion energy. The incident ions cause recoils within the system. The probability that a Se atom is recoiled is higher in the because of the abundance of Se-Se chains, when compared to the other two compositions, as well as their lower surface binding energy compared to Ge atoms, which are mostly found in four-fold coordinated tetrahedral structures. The recoiled Ge atoms have less kinetic energy than Se atoms while moving through the material, due to their higher lattice binding energy, which allows the Se atoms to migrate deeper into the material through collision cascades. Due to the low energy of the incident ions, many recoiled atoms lose their energy before sputtering an atom from the target. Therefore, the largest decrease in the Ge content is exhibited in these films since more Se atoms are recoiled deeper into the films, resulting in more Ge atoms being sputtered from the surface. The films a minimal change in composition, except when exposed to ions with the highest energy. The change in composition is not as exaggerated as the due to the reduction of Se-Se chains, which is also evident in the lack of compositional changes in the films, as verified in the Raman spectra described in detail later. The EDS data also revealed a presence of interstitial Ar atoms in the film. Statistically the number of absorbed is higher in the films, diminishes in the films and almost vanishes in the composition. We suggest that the accommodation of into the chalcogenide matrix depends upon the network packing fraction, since we noticed a reand films, duced presence of into the

2859

which have the highest packing fraction when compared to the films–Fig. 7. The introduction of ions also affects the surface properties of the material related to the appearance of hillocks at the highest ion fluence. Hillocks are small mounds or hills that are detected on the film surface. Greater number and/or larger hillocks will cause different film thickness in close proximity, which can alter the device performance from one write cycle to the next. When we connect the introduction of ions and appearance of hillocks with the results presented in Fig. 6., the decrease in the hillock height for the glasses at the highest ion energy confirms the role of interstitial Ar contributing to the changes in surface morphology. This effect is not as expressed for the other compositions where there is a limited amount of incorporated within the chalcogenide film matrix. Comparing this data to the Raman structural information in the films, it reveals that at the lowest energy (150 eV), structural reorganization predominantly occurs since the increased material loss with ion flux is coupled with the decrease of the Se-Se bonding. This means that the structure is stabilizing after the interaction with ions through three-dimensional self-organization. Self-organization is an important factor in disordered systems [23]. In the particular case, we suggest the decrease of Se chain vibrations coupled with decrease of Ge in the film composition is related to the formation of 3D tetrahedral structure (CS or ES) due to the breakage of some Ge-Ge bonds, leaving free Ge valent states to react with Se. Such structural variation has also been found with ion interaction in the Ge-Te system [24]. This tendency is sustained by the irradiation with 300 eV ions up to particles cm , after which an increase in the Se-Se bonding was manifested. We relate the last fact to the substantial loss of Ge at the highest ion fluence. In difference from [10], we obtained a red shift of the Se-Se Raman modes–Fig. 9. due to the loss of Ge in the films. These modes are very composition sensitive [25] and we attribute this as the reason behind the development of the peak shift at lower ion fluence. At higher ion fluences, due to increased accumulation of into the chalcogen matrix, acting as an external pressure over the chalcogenide network, the observed shift in the Se-Se Raman mode at the lowest fluence was sustained up to the highest fluence regardless of the ion energy. Under these radiation conditions, this film behaves like one with composition similar to Ge-rich compositions, even though the film incurred a loss of Ge atoms due to sputtering. In the films containing 33 at.% and 40 at.% Ge, the weakest bond in this system is one connecting two neighboring Ge atoms which are part of the ETH-like units. Ion irradiation produces changes in these bonds as shown in Fig. 4(b). and (c). For the composition, the saturation of all bonds and lack of dangling bond defects, because of its stoichiometric composition ( ), is an important factor for the reduced compositional variations demonstrated in Fig. 3(b). Similarly, the minimal changes to the areal intensity of the Ge-Ge bond are also attributed to the near stoichiometric composition of these films. With increasing amount of Ge, as in the films, we suggest an effect similar to the swelling reported for amorphous Ge films [26] which occurs in order to reduce the packing density in the system, since it is very high for this composition [20]. This high compactness of the structure significantly increases

2860


Fig. 8. IV curves in different cells of the 9th row for the fabricated RCBM array. (a)

Fig. 9. Peak locations of the stretch of the Se-Se Raman active (a) Fluence ( Particle cm ). (b) Fluence ( Particle cm ).

device. (b)

stretch combining the contributions of

the stress within the films. It is for this reason that the areal intensity of the Ge-Ge bond Raman vibrations increases, and also the reduction of packing through the formation of CS structures. The TRIM simulations suggest that the ion-film interaction occurring on the surface, which should affect the surface roughness of the films since the highest total ionization of electrons occurs in this region. The surface roughness of the chalcogenide films is an important factor for the performance of the RCBM devices since it corresponds to the effective distance for the silver ions building the conductive bridge. The results of the AFM studies show that the sputtering of individual atoms does change the surface morphology. Additionally, the restructuring due to relaxation of the network to reduce stress caused by interaction with the ions, also contributes to the observed surface changes. The material characterization data shows that the Ge richest composition ( ) presents the best stability, and the electrical performance characteristics of these devices formed by ion bombardment confirm this observation. Overall, as presented in Fig. 8(b).and (c), the devices based on composition containing over 33 at.% germanium exhibit good stability, and their performance is similar to those reported for devices obtained by conventional lithographic methods [27]. We attribute the similarity to the fact that incorporated are not electronically active dopants affecting the electrical performance of the devices, and the surface damage caused by the ions does not

device. (c)

rings and Se chains in

Device.

Raman spectra.

critically affect the device performance. This data supports the viability of this method for the fabrication of memory device (memristors) arrays. V. CONCLUSION inIn this paper, we present a study of low energy teraction with chalcogenide glass thin films. Our data of irradiated chalcogenide glasses, ranging from chalcogen-rich to chalcogen-poor, demonstrate that the chalcogen-richest glasses are most sensitive to Ar ion irradiation, towards structural and compositional changes. These effects can be related to the extraction of Ge atoms out of the chalcogenide matrix during irradiation, and structural reorganization to accommodate the stress caused by the introduction of Ar. An increase of the Ge content in the films leads to higher compactness and rigidity of the structure. This data was confirmed with AFM imaging and TRIM simulation. The energies applied are tailored for the needs of device formation and performance due to their relatively low effect over the chalcogenide matrix, and are representative of possible ambient space conditions that devices could experience while under cover from shielding. At the same time, these energies provide a well-established sputtering velocity, yielding defined RCBM device arrays with excellent and switching device characteristics, such as uniform behavior. Integrated circuits based solely on memristor arrays


show a great potential for new circuit solutions as demonstrated by Cali et al. [28]. ACKNOWLEDGMENT The authors would like to thank J. Reed of DTRA for his support. They would also like to recognize the Surface Science lab at Boise State University for AFM use and P. Davis for his assistance in performing AFM studies. REFERENCES [1] A. Zakery and S. Elliott, “Optical properties and applications of chalcogenide glasses: A review,” J. Non-Cryst. Sol., vol. 330, pp. 1–12, 2003. [2] J. Sanghera and I. Aggarwal, “Active and passive chalcogenide glass optical fibers for IR applications: A review,” J. Non-Cryst. Sol., vol. 256, pp. 6–16, 1999. [3] Jursa, Handbook of geophysics and the space environment, A. S. Jursa, Ed. Springfield, VA, USA: Air Force Geophysics Laboratory, 1985. [4] J. P. Gardner, J. C. Mather, M. Clampin, R. Doyon, M. A. Greenhouse, H. B. Hammel, J. B. Hutchings, P. Jakobsen, S. J. Lilly, and K. S. Long, “The James Webb Space Telescope,” Space Sci. Rev., vol. 123, pp. 485–606, 2006. [5] T. Tsvetkova, S. Balabanov, B. Amov, A. Djakov, and I. Wilson, “Surface morphology changes in ion implanted chalcogenide films after annealing,” Nucl. Instrum. Meth. Phys. Res., Sect. B, vol. 80, pp. 1264–1267, 1993. [6] I. Ivan, S. Szegedi, L. Daroczi, I. Szabo, and S. Kokenyesi, “Deuteron irradiation induced changes in amorphous AsSe films,” Nucl. Instrum. Meth. Phys. Res. Sect. B: Beam Interactions with Materials and Atoms, vol. 229, pp. 240–245, 2005. [7] S. Kokenyesi, I. Iván, V. Takáts, J. Pálinkás, S. Biri, and I. Szabo, “Formation of surface structures on amorphous chalcogenide films,” J. Non-Cryst. Sol., vol. 353, pp. 1470–1473, 2007. [8] E. H. Lee, “Ion-beam modification of polymeric materials–fundamental principles and applications,” Nucl. Instrum. Meth. Phys. Res., Sect. B, vol. 151, pp. 29–41, 1999. [9] M. S. Ailavajhala, Y. Gonzalez-Velo, C. Poweleit, H. Barnaby, M. N. Kozicki, K. Holbert, D. P. Butt, and M. Mitkova, “Gamma radiation induced effects in floppy and rigid Ge-containing chalcogenide thin films,” J. Appl. Phys., vol. 115, pp. 043502-1–9, 2014. [10] M. S. Kamboj, G. Kaur, R. Thangaraj, and D. Avasthi, “Effect of heavy ion irradiation on the electrical and optical properties of amorphous chalcogenide thin films,” J. Phys. D: Appl. Phys., vol. 35, p. 477, 2002. [11] N. F. Mott and E. A. Davis, Electronic Processes in Non-Crystalline Materials. London, U.K.: Oxford Univ. Press, 2012. [12] M. F. Thorpe, “Networks, flexibility and mobility in,” in Encyclopedia of Complexity and Systems Science. New York, NY, USA: Springer, 2009, pp. 6013–6024.

2861

[13] M. R. Latif, T. L. Nichol, M. Mitkova, D. A. Tenne, I. Csarnovics, S. Kokenyesi, and A. Csik, “Ion beam effect on Ge-Se chalcogenide glass films: Non-volatile memory array formation, structural changes and device performance,” in Proc. IEEE Workshop Microelectronics And Electron Devices (WMED), 2014, pp. 1–4. [14] Z. Borisova, Glassy semiconductors. New York, NY, USA: Plenum, 1981. [15] Y. Nedeva, T. Petkovaa, E. Mytilineoub, and P. Petkov, “Compositional dependence of the optical properties of the Ge-Se-Ga glasses,” J. Optoelectron. Adv. Mater., vol. 3, pp. 433–436, 2001. [16] M. Jin, P. Boolchand, and M. Mitkova, “Heterogeneity of molecular thin films,” J. Non-Cryst. structure of Ag photo-diffused Sol., vol. 354, pp. 2719–2723, 2008. [17] G. H. Ivanov, B. T. Kolomiets, V. M. Lyubin, and V. P. Shilo, in Proc. Intern. Conf. Amorph. Semic., 1972, vol. 1, p. 88. [18] P. M. Bridenbaugh, G. P. Espinosa, J. E. Griffiths, J. C. Phillips, and J. Raman line in P. Remeika, “Microscopic origin of the companion ,” Phys. Rev. B, vol. 20, pp. 4140–4144, 1979. glassy [19] K. Jackson, A. Briley, S. Grossman, D. Porezag, and M. Pederson, and aA first-principles “Raman-active modes of astudy,” Phys. Rev. B, vol. 60, pp. R14985–R14989, 1999. [20] V. Georgieva, M. Mitkova, P. Chen, D. Tenne, K. Wolf, and V. Gadgas sorption studies of films using quartz janova, “ crystal microbalance,” Mater. Chem. Phys., vol. 137, pp. 552–557, 2012. [21] A. Feltz, Amorphous Inorganic Materials and Glasses. Weinheim, Germany: VCH, 1993. [22] P. Chen, M. Mitkova, D. A. Tenne, K. Wolf, V. Georgieva, and L. Thin Films for Vergov, “Study of the sorption properties of gas sensing,” Thin Solid Films, vol. 525, pp. 141–147, 2012. [23] P. Boolchand, G. Lucovsky, J. C. Phillips, and M. F. Thorpe, “Selforganization and the physics of glassy networks,” Philosophical Mag., vol. 85, pp. 3823–3838, Nov 2005. [24] R. De Bastiani, E. Carria, S. Gibilisco, M. Grimaldi, A. Pennisi, A. Gotti, A. Pirovano, R. Bez, and E. Rimini, “Ion-irradiation-induced selective bond rearrangements in amorphous GeTe thin films,” Phys. Rev. B, vol. 80, p. 245205, 2009. [25] X. Feng, W. J. Bresser, and P. Boolchand, “Direct evidence for stiffness threshold in chalcogenide glasses,” Phys. Rev. Let., vol. 78, pp. 4422–4425, 1997. [26] S. Mayr and R. Averback, “Ion-irradiation-induced stresses and swelling in amorphous Ge thin films,” Phys. Rev. B, vol. 71, p. 134102, 2005. [27] Y. Gonzalez-Velo, H. Barnaby, M. Kozicki, P. Dandamudi, A. Chandran, K. Holbert, M. Mitkova, and M. Ailavajhala, “Total-ionizingdose effects on the resistance switching characteristics of chalcogenide programmable metallization cells,” IEEE Trans. Nucl. Sci., vol. 60, no. 6, pp. 4563–4569, 2013. [28] E. Gale, B. de Lacy Costello, and A. Adamatzky, “Emergent spiking in non-ideal memristor networks,” Microelectron. J., pp. 1–15, 2014.