Chalcogenide thin films that undergo reversible phase changes show promise for use in next-generation nanophotonics, microelectronics, and other emerging technologies. One of the many studied compounds, Ge2Sb2Te5, has demonstrated several useful properties and performance characteristics. However, the efficacy of benchmark Ge2Sb2Te5 is restricted by amorphous phase thermal stability below ∼150 °C, limiting its potential use in high-temperature applications. In response, previous studies have added a fourth species (e.g., C) to sputter-deposited Ge2Sb2Te5, demonstrating improved thermal stability. Our current research confirms reported thermal stability enhancements and assesses the effects of carbon on crystalline phase radiation response. Through in situ transmission electron microscope irradiation studies, we examine the effect of C addition on the amorphization behavior of initially cubic and trigonal polycrystalline films irradiated using 2.8 MeV Au to various doses up to 1 × 1015 cm−2. It was found that increased C content reduces radiation tolerance of both cubic and trigonal phases.

Phase change materials offer unique opportunities for emerging technologies such as metasurfaces, nanophotonics, and microelectronics. In particular, chalcogenide thin films can have their structures set through transitions, typically utilizing heat, that involves rapid switching between states. These structural transitions also lead to property differences, such as electronic resistivity or optical reflectance, which enable their use in different applications including phase-change memory devices.1 

The more commonly studied phase change thin film is Ge2Sb2Te5 (GST). GST can be switched rapidly between amorphous and crystalline states through changes in temperature, and the material exhibits large, corresponding changes in key properties.1 Crystalline GST forms the metastable, rock-salt cubic ( F m 3 ¯ m ) structure when heated to between ∼125 and ∼170 °C, depending on the heating rate.2 The stable trigonal structure ( P 3 ¯ m 1 ) is formed above ∼250 °C.3 GST has been identified as a rapidly transitioned material exhibiting decananosecond (10−8 s) switch times.4 For use in high-temperature applications, increased amorphous phase stability is needed. Dopants can be added to modify the crystallization temperature. Toward this end, adding Se,5 N,6,7 O,7 Al,8 Cu,8 Ag,9 In,10 and C11,12 have increased the thermal stability of amorphous GST and related compounds. C-doped GST has exhibited reduced power consumption and enhanced endurance in prior studies.13,14

Despite interest in these materials,15 relatively few reports have explored the effects of radiation on chalcogenide films. Ion irradiation has been used to cause amorphization, crystallize amorphous films, and tune the optical and electrical properties. Most of these studies have explored the fundamental effects of radiation on amorphous GST.16–23 For example, heavy ion irradiation with 120 keV Sb was used to modify the atomic structure of amorphous films so to enhance the crystallization kinetics compared with non-irradiated material.19 A separate, high energy 120 MeV Ag irradiation study has also demonstrated an ability to crystallize amorphous GST,18 leading to an awareness regarding potential ion-induced heating effects at high doses. The response of crystalline GST to ion irradiation has received less attention, but recent studies have revealed important details. Irradiation with 150 and 180 keV Ar induced amorphization at high doses with different threshold doses determined for both cubic and trigonal phases.16,17,23 Interestingly, irradiation of trigonal GST leads to the formation of the cubic phase prior to final, full amorphization. Despite these recent observations, relatively little is known about the radiation response of quaternary alloys of GST. Specifically, it is not known whether the addition of C (or other fourth species) bolsters or reduces radiation tolerance compared with undoped GST.

In this work, we use in situ transmission electron microscopy (TEM) to investigate the radiation stability of undoped and C-doped GST. In situ TEM offers the ability to observe changes in material microstructure, defects, and phase in real time, capturing the transient processes during applied external stimuli including high energy ion bombardment. In situ TEM generally offers a view of materials evolution during heating, cooling, mechanical deformation, or irradiation through electron beam imaging, electron diffraction, or both. TEM measurements are well-suited for studies of thin films due to the elastic scattering of electrons within the thin film material. Also, the amorphous to crystalline phase transformation temperatures of GST and GST with low carbon addition are accessible using in situ TEM heating stages.24,25 Recent in situ TEM studies have elucidated the nucleation and growth of crystalline phases within GST and other chalcogenides using advanced heating stages and high-speed, low-electron dose imaging to minimize electron beam effects,24–28 as electron beam irradiation of GST has been shown to both crystallize and amorphize GST.24,29

To be clear, in situ TEM with concurrent high energy ion irradiation is utilized for this study. This method is a well-established technique often utilized for studies in the nuclear materials community with several facilities having been developed for these studies.30,31 In the current study, we use electrons to monitor the phase and microstructure subjected to heavy ion bombardment. In situ characterization is advantageous, providing an opportunity to view the onset of phase change to full conversion of a material. This method of in situ, ion irradiation TEM provides an opportunity to systematically evaluate samples of different C concentrations to identify differences in radiation response.

Thin films of (Ge2Sb2Te5)1−xCx (0 ≤ x ≤ 0.12) were synthesized using direct current (DC) magnetron sputter deposition using a Unifilm Co. PVD-300 system and ultrahigh purity Ar background. Films were 90 nm-thick and deposited on top of 30 nm-thick, amorphous SiN (SPI Inc.) windows for plan-view TEM analysis. Specimens in the as-deposited condition were confirmed amorphous. C content was quantified via both separate wavelength dispersive spectroscopy (WDS) experiments utilizing thicker films and via scanning transmission electron microscopy elemental x-ray dispersive spectroscopy (STEM-EDS) on the same films used in this study.32 

The in situ ion irradiation transmission electron microscope (I3TEM) at Sandia National Laboratories was used to observe an irradiation-induced phase change.31 Samples were first annealed using one of two methods. Cubic-phase films were produced via in situ heating at a rate of 10 °C/min with full crystallization indicated by electron diffraction. Samples of 0, 2, 4, and 6% C required heating to 200 °C for 5 min whereas 12% C samples required a 450 °C, 5-min bake to establish the cubic phase. Control experiments involving EDS indicated no change in composition for these particular films. Cubic samples heated in situ were cooled for ∼30 min prior to ion irradiation. Experiments targeting the trigonal phase required higher temperatures. Annealing of 0, 2, and 4% C thin films to requisite 400 °C in vacuum led to substantial changes in composition as evidenced by EDS. Thus, an alternative approach was developed and implemented to preserve film composition when attaining the trigonal phase. Samples were alternatively baked to 400 °C for 15 min in ∼630 Torr of ultrahigh purity Ar in a separate vacuum furnace. The Ar overpressure suppressed Te outgassing and inhibited oxidation to preserve composition. The trigonal phase was confirmed by selected area electron diffraction (SAED) in samples having 0, 2, and 4% C used in subsequent irradiation experiments. Samples with greater amounts of C were not pursued for trigonal phase studies due to even higher temperature requirements.

These specimens were then irradiated in situ using a 6 MV Tandem accelerator with a 2.8 MeV Au4+ ion beam. Upstream measurements at a Faraday cup positioned prior to the TEM entry port indicated a stable beam flux in the range of 1 × 1011–1 × 1012 ions/cm²⋅s and associated current density between 7 × 10−8 and 6 × 10−7 A/cm². Doses reported throughout this paper, including threshold values for amorphization, are specified for the mounted sample position with effects of tilt angle included. Samples were nominally at room temperature when beginning irradiation. 2.8 MeV Au ions were chosen for reliable and stable ion fluxes and continue to investigate GST amorphization processes under ion beam bombardment in the nuclear collision regime to build off the work in Refs. 16, 17, and 22 with a similar range of nuclear and electronic stopping powers. Ions impacted the film nominally normal to the electron beam. The sample was tilted 30° to the electron beam (60° to the ion beam), as seen in Fig. 1(a). Bright-field TEM images and SAED patterns were obtained using a JEOL 2100 TEM operated at 200 keV.

FIG. 1.

(a) Experimental geometry showing GST film on SiN window supported on a Si substrate with impinging simultaneous electron and ion beams. The sample is tilted 30° relative to the electron beam, the ion beam is normal to the electron beam, and the ion beam hits the sample at 60° incidence. (b) SRIM-predicted damage (black) and Au implantation (red) profiles for 2.8 MeV Au4+ ions into an undoped cubic GST substrate at 60° incidence for a total dose of 1 × 1014 cm−2.

FIG. 1.

(a) Experimental geometry showing GST film on SiN window supported on a Si substrate with impinging simultaneous electron and ion beams. The sample is tilted 30° relative to the electron beam, the ion beam is normal to the electron beam, and the ion beam hits the sample at 60° incidence. (b) SRIM-predicted damage (black) and Au implantation (red) profiles for 2.8 MeV Au4+ ions into an undoped cubic GST substrate at 60° incidence for a total dose of 1 × 1014 cm−2.

Close modal

Damage thresholds, expressed as displacements per atom (dpa), are calculated with the stopping and range of ions in matter (SRIM) models where film density is estimated using the rule of mixtures between GST (either 4.086 × 1022 at/cm³ for cubic or 3.375 × 1022 at/cm³ for trigonal) and graphite (1.134 × 1023 at/cm³) using displacement threshold energies of 15 eV for Ge and 25 eV for Sb, Te, and C.33 Representative damage and Au implantation profiles are shown in Fig. 1(b) for an undoped cubic GST sample irradiated to a simulated dose of 1 × 1014 cm−2, indicating a relatively uniform damage profile throughout the sample thickness. The average dpa level throughout the sample in the undoped GST sample is 3.1 dpa for the simulated dose and decreases to 3.09, 3.0, and 2.9 dpa for the 2, 4, and 6% C-doped samples. For the undoped, trigonal GST sample, the average damage level throughout the sample depth is 2.9 dpa for a 2.8 MeV, 1 × 1014 cm−2 irradiation. Importantly, the implanted Au concentration is small for the experiments as intended. The maximum concentration of implanted Au is ∼0.005 at. % for the example shown, as the projected range of 2.8 MeV in GST at 60° incidence is ∼200 nm. The vast majority of incident Au ions traverse films and exit the SiN membranes as planned.

Image analysis is performed with ImageJ34 to assess phase change. Determination of the crystalline and amorphous fractions of cubic crystallized samples was completed by examining the average intensity of the (220) diffraction ring over time. As the irradiation progressed and as seen in the in situ video taken of the diffraction patterns, the intensity of all discrete crystalline diffraction rings decreases until there are just the broad, diffuse amorphous diffraction rings. The stack of images making up an irradiation was uploaded into ImageJ, an annular region of interest around the (220) diffraction ring was drawn, and the brightness and contrast normalized for all irradiation videos using the HiLo function. The average intensity of pixels within the annular region was then determined for each image slice in the stack, which was later converted to fluence step knowing the final amorphization fluence. The image processing scheme was repeated on an unirradiated sample that maintained the same crystallinity, and this processing scheme showed no change in crystallinity from the initial to final image slice, ensuring the accuracy of the process.

The as-deposited films are shown in Fig. 2, with associated electron diffraction patterns indicating an amorphous phase in all films, regardless of C content, prior to irradiation. The amorphous structure is indicated by broad, diffuse rings in the SAED pattern, which will be superimposed on that of the amorphous SiN windows.

FIG. 2.

TEM images and the corresponding SAED patterns of (a) 0, (b) 2, (c) 4, (d) 6, and (e) 12% C GST samples as deposited [(a)–(e)], after annealing to form cubic [(a′)–(e′)], and after irradiation [(a″)–(e″)], showing the progression from initially amorphous to crystalline to finally amorphous following irradiation.

FIG. 2.

TEM images and the corresponding SAED patterns of (a) 0, (b) 2, (c) 4, (d) 6, and (e) 12% C GST samples as deposited [(a)–(e)], after annealing to form cubic [(a′)–(e′)], and after irradiation [(a″)–(e″)], showing the progression from initially amorphous to crystalline to finally amorphous following irradiation.

Close modal

The microstructures and associated diffraction patterns following thermal treatments are also included in the second column of Fig. 2. To reiterate, the polycrystalline samples shown in Fig. 2 were baked inside the vacuum envelope of the TEM. The films were determined to be cubic exhibiting multiple, sharp, non-uniform diffraction rings indicating a nanocrystalline sample. A similar data set is shown in Fig. 3 for initially amorphous films converted to trigonal phase via ex situ baking. Diffraction patterns were indexed to confirm crystalline phases according to the structures described by Matsunaga.35,36 Example indexed SAED patterns are included in Fig. 4.

FIG. 3.

TEM images and the corresponding SAED patterns of trigonal (a) 0 C, (b) 2 C, and (c) 4% C GST samples after annealing [(a)–(c)] and after irradiation [(a′)–(c′)], showing the progression from crystalline to amorphous following irradiation.

FIG. 3.

TEM images and the corresponding SAED patterns of trigonal (a) 0 C, (b) 2 C, and (c) 4% C GST samples after annealing [(a)–(c)] and after irradiation [(a′)–(c′)], showing the progression from crystalline to amorphous following irradiation.

Close modal
FIG. 4.

Indexed SAED patterns for (a) cubic and (b) trigonal phases prior to irradiation.

FIG. 4.

Indexed SAED patterns for (a) cubic and (b) trigonal phases prior to irradiation.

Close modal

Figure 5(a) plots the grain size of crystallized films prior to irradiation. Grain sizes are determined from bright-field TEM images with one standard deviation of grains indicated. Grain size in the annealed state is reduced with increasing C content for the cubic phase. Refinement is evident as an eightfold reduction is generally observed when transitioning from x = 0 to 0.12 for the cubic phase. However, for the films crystallized into the trigonal phase, the grain size slightly increases as the C doping level increases from 0 to 4 at.% C. The high annealing temperature likely introduces a high enough thermal driving force to overcome the C doping resulting in grains of larger average size.

FIG. 5.

(a) Grain size and (b) amorphization dose as a function of carbon doping level for phase-pure cubic and phase-pure trigonal films. Error bars represent one standard deviation.

FIG. 5.

(a) Grain size and (b) amorphization dose as a function of carbon doping level for phase-pure cubic and phase-pure trigonal films. Error bars represent one standard deviation.

Close modal

The 2.8 MeV Au ion irradiation induced amorphization in all crystallized films. Figure 2 shows the same cubic films that were crystallized with an added third column showing the disordering resulting in amorphization from the ion irradiation. Similarly, Fig. 3 shows evidence of ion-induced amorphization of GST trigonal films having 0, 2, and 4% C. Diffraction patterns are included in each case with evidence of amorphous films associated with the relatively small number of diffuse, broad diffraction rings. Figure 5(b) shows the damage level requirements for amorphizing films. Additional details of these experiments are specified in Tables I and II. Note the damage levels and doses reported in figures and tables should be considered those required to fully amorphize films.

TABLE I.

Crystallization temperatures, crystallized grain size, damage to complete amorphization, and amorphization dose for all cubic samples.

Crystallization temperature (°C)Grain size (nm)Amorphization dose (ions/cm²)Damage to complete amorphization (dpa)
GST 133 ± 1.4 85 ± 46 1.07 × 1013 ± 4.93 × 1012 0.315 ± 0.054 
GST 2% C 157 ± 2.0 32 ± 19 3.04 × 1012 ± 5.10 × 1011 0.091 ± 0.025 
GST 4% C 155 ± 3.2 25 ± 16 1.61 × 1012 ± 4.64 × 1011 0.048 ± 0.015 
GST 6% C 156 ± 5.6 17 ± 6 6.52 × 1011 ± 2.18 × 1011 0.0189 ± 0.012 
GST 12% C 389 ± 42 10 ± 3 6.42 × 1011 ± 2.19 × 1011 0.018 ± 0.01 
Crystallization temperature (°C)Grain size (nm)Amorphization dose (ions/cm²)Damage to complete amorphization (dpa)
GST 133 ± 1.4 85 ± 46 1.07 × 1013 ± 4.93 × 1012 0.315 ± 0.054 
GST 2% C 157 ± 2.0 32 ± 19 3.04 × 1012 ± 5.10 × 1011 0.091 ± 0.025 
GST 4% C 155 ± 3.2 25 ± 16 1.61 × 1012 ± 4.64 × 1011 0.048 ± 0.015 
GST 6% C 156 ± 5.6 17 ± 6 6.52 × 1011 ± 2.18 × 1011 0.0189 ± 0.012 
GST 12% C 389 ± 42 10 ± 3 6.42 × 1011 ± 2.19 × 1011 0.018 ± 0.01 
TABLE II.

Crystallized grain size, damage dose to complete amorphization, and amorphization dose for all trigonal samples.

Anneal temperature (°C)Grain size (nm)Amorphization dose (ions/cm²)Dose to complete amorphization (dpa)
GST 400 221 ± 100 1.3 × 1014 ± 5 × 1012 4.2 ± 1 
GST 2% C 400 293 ± 71 2.8 × 1013 ± 5 × 1011 0.8 ± .08 
GST 4% C 400 325 ± 65 6.5 × 1012 ± 1 × 1011 0.2 ± .03 
Anneal temperature (°C)Grain size (nm)Amorphization dose (ions/cm²)Dose to complete amorphization (dpa)
GST 400 221 ± 100 1.3 × 1014 ± 5 × 1012 4.2 ± 1 
GST 2% C 400 293 ± 71 2.8 × 1013 ± 5 × 1011 0.8 ± .08 
GST 4% C 400 325 ± 65 6.5 × 1012 ± 1 × 1011 0.2 ± .03 

We note that electron-beam-only irradiation, primarily when focused to a point, can result in crystallization of the amorphous GST. Also, previous work by Jiang et al. demonstrates that extensive, focused electron beam irradiation can lead to progressive amorphization of initially crystalline GST.25 Initial work indicated electron beam-induced amorphization and crystallization in separate control experiments in bright-field imaging conditions. Therefore, in order to minimize electron-beam effects and isolate ion-irradiation processes, the TEM studies presented here were carried out in diffraction mode with the electron beam spread completely to minimize the electron dose during ion irradiation. We found no evidence of electron beam-induced amorphization or crystallization when observing in diffraction mode, conducted in TEM mode (not STEM mode) with a LaB6 filament, decreasing the electron beam flux on each sample.

In this work, it is observed that C addition decreases the cubic-crystallized GST grain size, mirroring a result observed by Song et al.14 It is proposed that C acts as a drag on GST limiting the diffusion and subsequent grain growth during annealing.14 The C doping has been shown to reduce GST grain size,14 attributed to reducing the diffusion of GST limiting grain growth during annealing. Also, we find that C raises the crystallization temperature of GST consistent with the previous work.12 We also observe higher temperatures needed to transform the cubic GST samples to the trigonal phase, similar to Ref. 12, which leads to a larger grain size and a greater spread in grain sizes (compared with the cubic phase).

We observe that 2.8 MeV Au irradiation amorphizes GST and the various C-doped GST films studied. The in situ amorphization doses measured for pure GST compare reasonably well to those reported previously giving credence to measured differences in films having carbon. For example, Hafermann et al.17 reports complete amorphization of pure GST when ndpa exceeds 0.6 (for cubic phases) and 3 (for hexagonal/trigonal phase) noting good agreement with earlier work by Privitera et al.,16 which utilized 180 keV Ar. Similarly, we find a higher damage level for complete amorphization of trigonal GST (ndpa = 4.2) compared with that of cubic (ndpa = 0.3). Differences between our experiments and past work may be a result of moderate sample heating in the TEM vacuum environment, variations in ion fluxes, the effects of more energetic (MeV) beams, or influence of combinations thereof.

The threshold doses for amorphization vary with carbon content for both trigonal and cubic phases, as shown in Fig. 5. Generally, the required dose to amorphization is reduced when adding carbon. For example, the amorphization dose is reduced by 1.5 decades for cubic GST 12% C compared with undoped, cubic GST. Accounting for density differences in films produces a similar variation in requirements for amorphization, expressed as dpa and ions fluence in Tables I and II. It was observed that larger doses are required to amorphize trigonal phase (Ge2Sb2Te5)1−xCx compared with the cubic phase, when comparing films of equivalent carbon content.

Radiation-induced amorphization is observed in contrast to radiation-induced grain growth that is observed in many nanocrystalline material systems, attributed to thermal spike (e.g., thermal events) via enhanced grain boundary mobility near damage cascades.37,38 Referred to as metamictization in the geologic and mineralogic communities, radiation-induced amorphization has been reported for numerous materials.39–42 Prior work has shown the critical amorphization dose tends to increase with irradiation temperature, indicating that temperature causes defect migration that reduces the point defect concentration for the ion bombardment. Indeed, prior experiments involving the irradiation of GST at 77 and 300 K show a >10x greater amorphization threshold at the higher temperature, which is consistent with enabled diffusion and recombination.22 Thus, it is assumed that the ion irradiation simply introduces a high number of lattice point defects that over time accumulate to completely amorphize the material. Irradiation of GST likely caused progressive amorphization through the creation of point defects in the lattice that eventually leads to amorphization.

Theories describing the irradiation-induced amorphization of a crystalline material posit that the amorphous fraction accumulation depends on the crystallization temperature and cascade size.42 Gibbons43 and Carter44 explained the mechanism of amorphization as individual damage cascades that accumulate into complete amorphization. Heavy ions at low temperatures directly transform a region of the crystal into the amorphous phase, and the accumulation of amorphous zones leads to complete amorphization. Weber45 explained amorphization of non-metals via defect accumulation, direct-impact amorphization, and cascade overlap. There are competing recrystallization and amorphization processes occurring and if the amorphization induced by the ion beam damage exceeds the thermal term dominated by external heating or temperature rises due to thermal spikes, then complete amorphization occurs.

The amorphous fraction (f) in Eq. (1) of a material during ion bombardment can be expressed as a function of dose41 as
f = 1 1 A + ( 1 A ) exp ( 2 ( 1 A ) 50 D ) ,
(1)
where D is the normalized ion dose and A is the extent of crystallization.42 In this case, we adapt the formula for f and replace D with dpa calculated from SRIM because we do not have in-depth modeling for the size of cascades in the GST volume. This change results in an empirical model for the progression of amorphization of the GST films, opposed to an analytical description applicable to all in situ TEM amorphization studies.

Figure 6 shows the progression of amorphization as the ion irradiation progresses, plotting the crystalline fraction vs dose curves for cubic (Ge2Sb2Te5)1−xCx [0 ≤ x ≤ 0.12] samples in this study in plot (a), as well as GST in prior studies with in situ monitoring of bulk un-doped GST specimens in plot (b).16,17 The data from Refs. 16 and 17 are digitized from a plot of reflectance data, correlating maximum reflectivity with crystallinity. The crystal fraction, defined as 1 − f, determined via analysis of the intensity of the crystalline diffraction rings, is shown to decrease as the irradiation fluence increases. In Fig. 6, the crystal fraction is normalized to the peak intensity of the (220) diffraction ring in each sample.

FIG. 6.

(a) Crystalline fraction vs damage for irradiated cubic samples in this study. (b) Reflectivity curves vs damage used as a surrogate for amorphization curves digitized from data in Refs. 16 and 17, respectively.

FIG. 6.

(a) Crystalline fraction vs damage for irradiated cubic samples in this study. (b) Reflectivity curves vs damage used as a surrogate for amorphization curves digitized from data in Refs. 16 and 17, respectively.

Close modal

Compared to the work in Refs. 16, 17, and 22, the pure GST specimens investigated in this study with in situ TEM techniques preserve crystallinity to somewhat similar damage levels. Carbon-containing films amorphized at lower dose. Fitting the crystalline fraction vs damage level curves in Fig. 6, to the above equation for the amorphous fraction, we solve for A for each cubic sample. The calculated values of A are shown as insets in Figs. 6(a) and 6(b). The value “A” represents the extent of crystallization, and values closer to one indicate the material, at the given irradiation condition, preserves crystallinity longer. The crystalline fraction (1 − f) as a function of ion dose of the 0 and 2% C-doped cubic-crystallized samples mirrors the defect accumulation curve, while at 4 and 12% C, the curves more closely resemble direct-impact amorphization.45 Thus, in addition to having a lower final amorphization dose, C doping of GST accelerates the amorphization of GST.

There are multiple GST film variables to consider in their impacts on the radiation tolerance, including C doping, grain size, and phase. The C doping has been shown to reduce cubic-phase grain size, attributed to reducing the diffusion of GST limiting grain growth during annealing.11–14 Along the same vein, during ion irradiation the presence of C may be limiting diffusion of ion irradiation-induced point defects. These would lead to the faster accumulation of defects at higher C doping concentrations, leading to a lower amorphization dose and potentially contributing to the enhanced kinetics for amorphization at higher C concentrations. DFT simulations by Chu et al.11 showed that C dopants in GST affect the local ordering and Borisenko et al.46 showed they form carbon-rich clusters, slowing the diffusion of the constituent Ge, Sb, and Te. The impact of C doping on the radiation tolerance is observed in both cubic- and trigonal-crystallized samples, displaying the same trend. Thus, across GST grain size and phase, the increased C content limiting diffusional recombination of irradiation-induced knock-on defects will contribute to decreasing the damage dose needed for amorphization.

The grain size differences are most prominent for the cubic-crystallized samples, as the grain size decreases by 70% as the C content increases from 0% to 4%, as compared to 45% for the trigonal-crystallized samples. The increased density of grain boundaries themselves may contribute to amorphization. Grain boundaries have been shown to both enhance radiation tolerance47,48 and accelerate amorphization.49–52 The grain boundaries are defects with lower lattice order, which results in lower displacement thresholds than the surrounding matrix and subsequently increases damage at and near grain boundaries.53,54 Thus, increasing the cubic-phase grain boundary density could lower the amorphization threshold by having more damage occurring at the higher density of grain boundaries due to having more atoms with lower coordination at interfaces. However, the grain boundary density trend only holds for the cubic-crystallized samples. The trigonal phase samples do not show a similar decrease in grain size (by extension, grain boundary density), and a decreasing amorphization dose is exhibited by trigonal films with increasing carbon content. The grain size slightly increases as C is added to the trigonal-crystallized samples, although there is a large scatter. Therefore, we expect that the intrinsic properties of the films enhance the radiation tolerance more strongly than grain boundaries. Altogether, it is surmised that the grain boundary density does not account solely for the differences in amorphization dose in either cubic- or trigonal-crystallized samples.

In this study, we observe that trigonal phase GST has a greater radiation tolerance to amorphize, compared with cubic-phase films. The work by Privitera et al.22 showed that under 150 keV Ar irradiation, the ion irradiation induces a similar progressive amorphization to what is observed here. The increase in radiation tolerance of the trigonal structure is mirrored, attributed to a transition to rock-salt structure before complete amorphization. In this work, we observe an increased radiation tolerance of the trigonal phase. Mechanistically, Privitera et al.22 proposed that a vacancy disordering induced a filling-in of van der Waals gaps in the trigonal phase during irradiation, inducing a transition to cubic phase at 15% displaced atoms, followed by further dose needed for amorphization of the newly formed cubic phase. The trigonal phase structure proposed by Matsunaga et al. is composed of stacked atomic layers with adjacent Te planes.37 These Te planes are predicted to be bonded weakly by van der Waals forces resulting in separation gaps of 2.8 ± 0.2 Å, exceeding the average planar spacing in the trigonal structure ∼1.8 Å55 and as these are filled with displaced Ge and Sb atoms, a transition to cubic phase is encouraged.22 

For the 2.8 MeV Au ion bombardment studied here, there is also the consideration of electronic stopping. The above discussion concerned nuclear stopping and knock-on point defect damage accumulation leading to amorphization. The electronic excitation method for amorphizing a material through melting and quenching of amorphous regions in tracks may be occurring in GST.56 In this theory, the amorphous tracks accumulate across the entire thickness of the GST film, due to the 90 nm thickness of the film and the Bragg peak being beyond the thickness of the film. With the grain boundaries being less-ordered regions within a polycrystalline material, a reduced grain size will increase the fraction of the sample that is intrinsically disordered before irradiation begins. Therefore, when the irradiation begins and forms amorphous tracks, these amorphous tracks are more likely to intersect with a grain boundary. This further widens the amount of material that is amorphous. At higher doses and continual irradiation, the tracks overlap as subsequent ion strikes hit already amorphous regions. Conversely, a coarse-grained material will require more amorphous tracks, and correspondingly, ion fluence, to interact with grain boundaries and thus amorphize the material. However, for the trigonal-crystallized samples in this study, the percent change in the grain size between 0 and 4% C is decreased; therefore, the amorphous tracks that may occur would intersect a similar density of grain boundaries, indicating that C doping plays a larger role than grain size. Similarly, in the cubic samples, the C addition, which limits grain size through limiting diffusion, may be limiting the diffusion of irradiation-induced point defects, which would lead to the accumulation of irradiation damage leading to amorphization. However, in this work, we do not investigate nor see evidence of ion tracks, but the inelastic thermal spike model may be appropriate to describe the formation of the amorphous region.

This work, having been conducted in situ in a TEM, does combine both electron and ion beam irradiation, both of which have the potential to damage the film and elevate its temperature. However, given the ion and electron beam currents chosen here, the effect of beam heating is expected to be minimal. Within the individual damage cascade there is localized melting; however, on the macroscopic film scale we do not reach the melting point of the entire GST film during the in situ irradiation due to ion beam heating. In this study, the majority of the ion energy is deposited in the GST via nuclear stopping, which would contribute to local melting of the sample during individual ion strikes. Local melting under individual ion strikes estimates temperatures on the order of 5200 °C during the cascade; however, re-solidification of the material occurs on the picosecond timescale and individual cascades are nanometers in diameter.57 The effect of the damage accumulated from each individual ion strike-induced melting and re-solidification will re-arrange the GST microstructure in tandem with the knock-on damage deposited. However, if the ion beam-induced heating is sufficient to increase the global temperature of the film, a small increase in film temperature may increase the amorphization dose, although further work is needed to verify this.58 Atomistic modeling is currently under way to decipher the multitude of effects and mechanisms that lead to amorphization of the undoped and C-doped GST, including the location of the solute C atoms within GST grains and response to heavy ion bombardment. A full analysis should include the effects of crystal structure and vacancy content characteristic of the different phases. The amorphization under irradiation is not limited to the in situ TEM experimental geometry and can be translated to ex situ irradiations with further characterization. This work represents the first in situ ion irradiation TEM study of GST to the authors knowledge. The insight from these studies provides a clear avenue to explore thermal and radiation stability of GST of various dopants and microstructures with enhanced spatiotemporal fidelity.

Chalcogenide thin films such as GST have been proposed for emerging electronic and optical applications as their crystalline-amorphous switching enables reconfigurable states. However, the limited stability of the chalcogenide amorphous phase limits its effectiveness, especially in thermal environments. Through in situ TEM annealing and irradiation, we show that C doping from 2% to 12% of GST increases the crystallization temperature and reduces the crystallized, cubic-phase grain size. Subsequent in situ, 2.8 MeV Au ion irradiation amorphizes the crystallized samples. The measured amorphization dose decreases as the C content increases from 0% to 12% for both cubic and trigonal phase GST. The amorphization process of a crystalline sample is discussed in relation to the ion bombardment dose, C content, GST phase, and the grain size of crystallized samples to offer design pathways for enhanced thermal and irradiation stability. The reduced tolerance for ion irradiation-induced amorphization in doped GST films is attributed to intrinsic effects of carbon.

The authors would like to thank Dr. Michael Abere for helpful discussions. The authors would like to thank C. Sobczak for sample preparation and P. Lu for EDS. We acknowledge the Laboratory Directed Research and Development program for providing funding for this study. This work was performed, in part, at the Center for Integrated Nanotechnologies, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science. Sandia National Laboratories is a multimission laboratory managed and operated by National Technology & Engineering Solutions of Sandia, LLC, a wholly owned subsidiary of Honeywell International, Inc., for the U.S. DOE's National Nuclear Security Administration under Contract No. DE-NA-0003525. The views expressed in the article do not necessarily represent the views of the U.S. DOE or the United States Government.

The authors have no conflicts to disclose.

Eric Lang: Data curation (equal); Formal analysis (equal); Investigation (equal); Visualization (equal); Writing – original draft (equal). Trevor Clark: Data curation (equal); Formal analysis (equal); Investigation (equal); Visualization (equal); Writing – review & editing (equal). Ryan Schoell: Investigation (equal); Visualization (equal); Writing – review & editing (equal). Khalid Hattar: Project administration (equal); Resources (equal); Supervision (equal); Writing – review & editing (equal). David P. Adams: Funding acquisition (equal); Project administration (equal); Resources (equal); Supervision (equal); Writing – review & editing (equal).

The data that support the findings of this study are available from the corresponding authors upon reasonable request.

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