the XRD spectrum was typical of an amorphous phase and the electrical resistivity was comparable to that measured in the as deposited amorphous film. The complete amorphization was detected by X-ray diffraction in glancing angle, to enhance the signal from the thin phase change layer. XRD measurements have been carried out using a SIEMENS D5005 diffractometer operating with Cu K ( =0.15406 nm). The X-ray beam was incident at 0.7° on the sample surface and the detector angle was incremented with an angular step of 0.02°. Figure 2 shows in red the XRD spectra of 70 nm thick crystalline films, as deposited at 250°C, with typical rocksalt patterns, and, in blue, the same film after ion implantation of Ge. No XRD peak is visible anymore and the pattern is typical of amorphous materials.

Fig. 2. XRD spectra of Ge2Sb2Te5 deposited at 250°C, in the rocksalt structure (red), and after ion implantation with Ge at 130KeV, 9e13/cm2 (blue). After implantation the XRD peak is typical of an amorphous material.

Fig. 3. Sheet resistance of melt quenched crystalline films versus laser intensity. The higher sheet resistance is measured in completely amorphized films.

the XRD spectrum was typical of an amorphous phase and the electrical resistivity was comparable to that measured in the as deposited amorphous film. The complete amorphization was detected by X-ray diffraction in glancing angle, to enhance the signal from the thin phase change layer. XRD measurements have been carried out using a SIEMENS D5005 diffractometer operating with Cu K ( =0.15406 nm). The X-ray beam was incident at 0.7° on the sample surface and the detector angle was incremented with an angular step of 0.02°. Figure 2 shows in red the XRD spectra of 70 nm thick crystalline films, as deposited at 250°C, with typical rocksalt patterns, and, in blue, the same film after ion implantation of Ge.

Fig. 2. XRD spectra of Ge2Sb2Te5 deposited at 250°C, in the rocksalt structure (red), and after ion implantation with Ge at 130KeV, 9e13/cm2 (blue). After implantation the XRD peak is

Fig. 3. Sheet resistance of melt quenched crystalline films versus laser intensity. The higher

sheet resistance is measured in completely amorphized films.

typical of an amorphous material.

No XRD peak is visible anymore and the pattern is typical of amorphous materials.

Melt-quenched films were obtained by using a Ruby laser operating at 695nm. Samples were irradiated with a 30 ns laser pulse, employing a spot diameter of 3.5 mm. The energy density was varied and the effect of film irradiation was monitored by measuring the sheet resistance using a four point probe.

Figure 3 shows the sheet resistance as a function of laser intensity. Below 30 mJ/cm2 the measured sheet resistance is around 1000Ohm/square, typical of crystalline GST. This value gradually increases with laser intensity and reaches a saturation value of 5E7 Ohm/square for energy densities above 40mJ/cm2.

Data shown in fig. 3 can be explained by considering the depth of melt versus energy density. The heating and melting produced by the laser pulse can be described by considering the heat diffusion equation with the addition of a source term driven by the absorption of laser radiation (Baeri, 1982). If the film is spatially uniform, then away from the edges of the irradiated area any temperature variation will occur only along the axis perpendicular to the sample surface. This simplifies the problem to one dimension and it is described by the equation

$$\frac{\partial \mathbf{T}}{\partial \mathbf{t}} = \frac{a}{\rho \mathbf{C}\_p} \mathbf{I}(\mathbf{z}, \mathbf{t}) + \frac{1}{\rho \mathbf{C}\_p} \frac{\partial}{\partial \mathbf{z}} \left( \kappa \frac{\partial \mathbf{T}}{\partial \mathbf{z}} \right) \tag{1}$$

where T is the temperature at time t and depth z, is the absorption coefficient of the illuminated material, with specific heat capacity Cp and density , is the thermal conductivity and I is the laser power density.

$$\mathbf{I}(\mathbf{t}, \mathbf{z}) = \mathbf{I}\_0(\mathbf{t}) \text{ (1-R) } \exp(-\alpha \mathbf{z}) \tag{2}$$

For the same laser pulse duration, the melt region depth can be calculated by dividing the sample into slices of thickness z and discretising time into units t. The temperature of each slice is then calculated from the energy deposited into the material by absorption of laser light and heat transfer by diffusion. The parameters used for calculations are shown in table 1.

Table 1. Parameters used for numerical calculations of the depth of melting region under irradiation by a 30 ns laser pulse

The results of simulations are shown in figure 4, where the melt depth is plotted as a function of time for different energy densities. According to the measured sheet resistance no melting, which means no amorphization, is expected for energy density below 30mJ/cm2. For energy densities above this threshold value, the melting depth increases as the energy density increases.

thickness.

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 135

material. For melt quenching, samples irradiated at two different laser intensities have been tested. Both laser intensities values are high enough to melt and quench the entire film

Fig. 6. Conductivity versus temperature for as deposited (a), implanted (b) and laser

twice the activation energy for conduction (Mott & Davis, 1979).

The carrier transport in amorphous GST in the temperature range we studied (40-120 ºC), is dominated by the band conduction through extended states, in which the charge carriers hop between trap centers via excitations above the mobility edge. Therefore the conductivity of amorphous samples is exponentially dependent on the energy gap, whose value is about

annealed (c) amorphous GST.

Fig. 4. Melted depth versus time as simulated for several energy densities.

Figure 5 shows the simulated melt depth as a function of laser energy density. The arrows indicate the energy density at which the thickness of melted material is expected to reach the film thickness of 70 nm. The simulated value of 40 mJ/cm2 required to melt the entire film agrees very well with the energy density at which sheet resistance saturation is observed, indicating that all the material is converted into the amorphous phase by melt quenching.

Fig. 5. Energy density versus melt thickness. For 70 nm thick film complete amorphization is expected above 40mJ/cm2.

For comparison, amorphous films were also deposited at room temperature over SiO2 covered Si substrates, from a single target by using RF sputter. Thickness of amorphous films was 50 nm. The stability of the amorphous resistance has been studied through sheet resistance measurements at different temperatures, in the range 40-120°C.

Figure 6 (a), (b) and (c) shows as black symbols the conductivity versus the reciprocal of temperature for as deposited, implanted and melt quenched amorphous films, respectively. The figure also shows as red symbols the conductivity obtained after annealing at temperatures below 140°C, i.e. not so high to produce the conversion into the crystalline phase. Four different samples have been measured for amorphous and as implanted

Fig. 4. Melted depth versus time as simulated for several energy densities.

quenching.

expected above 40mJ/cm2.

Figure 5 shows the simulated melt depth as a function of laser energy density. The arrows indicate the energy density at which the thickness of melted material is expected to reach the film thickness of 70 nm. The simulated value of 40 mJ/cm2 required to melt the entire film agrees very well with the energy density at which sheet resistance saturation is observed, indicating that all the material is converted into the amorphous phase by melt

Fig. 5. Energy density versus melt thickness. For 70 nm thick film complete amorphization is

For comparison, amorphous films were also deposited at room temperature over SiO2 covered Si substrates, from a single target by using RF sputter. Thickness of amorphous films was 50 nm. The stability of the amorphous resistance has been studied through sheet

Figure 6 (a), (b) and (c) shows as black symbols the conductivity versus the reciprocal of temperature for as deposited, implanted and melt quenched amorphous films, respectively. The figure also shows as red symbols the conductivity obtained after annealing at temperatures below 140°C, i.e. not so high to produce the conversion into the crystalline phase. Four different samples have been measured for amorphous and as implanted

resistance measurements at different temperatures, in the range 40-120°C.

material. For melt quenching, samples irradiated at two different laser intensities have been tested. Both laser intensities values are high enough to melt and quench the entire film thickness.

Fig. 6. Conductivity versus temperature for as deposited (a), implanted (b) and laser annealed (c) amorphous GST.

The carrier transport in amorphous GST in the temperature range we studied (40-120 ºC), is dominated by the band conduction through extended states, in which the charge carriers hop between trap centers via excitations above the mobility edge. Therefore the conductivity of amorphous samples is exponentially dependent on the energy gap, whose value is about twice the activation energy for conduction (Mott & Davis, 1979).

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 137

Fig. 7. Conductivity versus reciprocal temperature for undoped amorphous GST (diamonds)

Figure 7 shows the conductivity of amorphous samples as a function of the reciprocal temperature, measured in undoped GST (black diamonds) and in samples with: 1.5 at.% of O (full squares), 6 at.% of O (open squares), 1.5 at.% of N (full triangles) 10 at.% of N (open triangles), 1.5 at.% of F (full circles) and 6 at.% of F (open circles). As shown by the arrows in Fig.7, in the case of nitrogen or oxygen doping, the conductivity decreases as the dopant concentration increases. In the case of GST doped with fluorine, instead, the conductivity

Fig. 8. Activation energy for conductivity as a function of doping concentration, obtained for

Figure 8 shows the activation energy versus doping concentration, as obtained by linearly

In samples doped with nitrogen or oxygen the activation energy for conduction slightly increases, while for fluorine doping the activation energy remains almost the same. Largest variations have been obtained in the case of nitrogen: by doping GST with 10 at.% of N, the conductivity decreases of one order of magnitude and E increases from 0.42 *±* 0.01 eV,

nitrogen (diamonds), oxygen (squares) and fluorine (triangles).

measured in undoped samples, up to 0.47 *±* 0.01 eV.

and samples doped with oxygen (squares), nitrogen (triangles) or fluorine (circles).

increases.

fitting data shown in Fig. 7.

In as deposited samples the mobility gap, measured by linearly fitting the data shown in Fig.6, is 0.86 eV and its value does not change after annealing at 130°C for 10 minutes. Therefore in as deposited GST we do not observe the resistance drift reported for PCM devices.

In films amorphized by ion implantation, instead, the conductivity decreases after annealing at 120°C. The mobility gap also changes from 0.7eV, obtained in fresh samples, to 0.82-0.84 eV, a value very close to the as deposited amorphous. A similar behavior is also observed for melt quenched amorphous films.

The drift of activation energy for conduction in amorphous GST has been already observed in amorphous cells and it has been explained considering two possible candidates: a change of the bandgap value due to mechanical strain; a change of the donor acceptor-like defect density (Pirovano et al., 2004). Under compressive strain the bandgap of several chalcogenide materials decreases by as much as 2eV/ relative volume variation (Kastner, 1972). It is known that the phase change cause a compression of 6%. Therefore meltquenched films could exhibit lower bandgap due to accumulated mechanical stress. This stress can be then released, producing and increase of E, thanks to the bandgap increase during stress release.

On the other hand, it has been observed that ion implantation can produce densification in phase change materials (Rimini et al., 2009), obtaining amorphous films of substantially higher mass density than as deposited amorphous films (Raoux et al., 2010). Therefore we can reasonably consider the observed modification of E in both melt quenched and ion implanted film as due to compressive stress release.

Another possible explanation is given by recent calculations (Akola et al., 2010) and involves the role of Ge atoms bonding. Atomistic simulations of molecular dynamics performed on amorphous phase-change materials have shown that in melt quenched and as deposited amorphous GST, Germanium atoms have different environments: predominantly tetrahedral in as deposited GST and distorted octahedral in melt quenched. EXAFS data (Kolobov et al., 2004) indicate that in amorphous GST Ge atoms are prevalently tetrahedrally coordinated, with considerably shorter bond lengths (2.61-2.63 Å) than in the crystal, where Ge atoms occupy the octahedral position with Ge-Te bond length 2.83 Å. Therefore the decrease of activation energy for conduction observed in films amorphized by ion implantation and in melt quenched GST could be determined by a larger fraction of Ge atoms with distorted octahedral coordination.

## **4. Doping by Ion beam: Amorphous properties**

Films doped with nitrogen, oxygen or fluorine have been obtained by ion implantation on amorphous Ge2Sb2Te5. The energy of the ions was 15 keV, in order to have distributions peaked at the center of the film. Different fluencies were implanted, obtaining peak concentrations of 1.5 at.% and 10 at.% of nitrogen; 1.5 at.% and 6 at.% of oxygen; 1.5 at.% and 6 at.% of fluorine.

The amorphous properties were studied by electrical measurements. The sheet resistance was measured *in situ* using a four-point probe configuration, with the employment of a thermal chuck.

In as deposited samples the mobility gap, measured by linearly fitting the data shown in Fig.6, is 0.86 eV and its value does not change after annealing at 130°C for 10 minutes. Therefore in as deposited GST we do not observe the resistance drift reported for PCM

In films amorphized by ion implantation, instead, the conductivity decreases after annealing at 120°C. The mobility gap also changes from 0.7eV, obtained in fresh samples, to 0.82-0.84 eV, a value very close to the as deposited amorphous. A similar behavior is also observed

The drift of activation energy for conduction in amorphous GST has been already observed in amorphous cells and it has been explained considering two possible candidates: a change of the bandgap value due to mechanical strain; a change of the donor acceptor-like defect density (Pirovano et al., 2004). Under compressive strain the bandgap of several chalcogenide materials decreases by as much as 2eV/ relative volume variation (Kastner, 1972). It is known that the phase change cause a compression of 6%. Therefore meltquenched films could exhibit lower bandgap due to accumulated mechanical stress. This stress can be then released, producing and increase of E, thanks to the bandgap increase

On the other hand, it has been observed that ion implantation can produce densification in phase change materials (Rimini et al., 2009), obtaining amorphous films of substantially higher mass density than as deposited amorphous films (Raoux et al., 2010). Therefore we can reasonably consider the observed modification of E in both melt quenched and ion

Another possible explanation is given by recent calculations (Akola et al., 2010) and involves the role of Ge atoms bonding. Atomistic simulations of molecular dynamics performed on amorphous phase-change materials have shown that in melt quenched and as deposited amorphous GST, Germanium atoms have different environments: predominantly tetrahedral in as deposited GST and distorted octahedral in melt quenched. EXAFS data (Kolobov et al., 2004) indicate that in amorphous GST Ge atoms are prevalently tetrahedrally coordinated, with considerably shorter bond lengths (2.61-2.63 Å) than in the crystal, where Ge atoms occupy the octahedral position with Ge-Te bond length 2.83 Å. Therefore the decrease of activation energy for conduction observed in films amorphized by ion implantation and in melt quenched GST could be determined by a larger fraction of Ge

Films doped with nitrogen, oxygen or fluorine have been obtained by ion implantation on amorphous Ge2Sb2Te5. The energy of the ions was 15 keV, in order to have distributions peaked at the center of the film. Different fluencies were implanted, obtaining peak concentrations of 1.5 at.% and 10 at.% of nitrogen; 1.5 at.% and 6 at.% of oxygen; 1.5 at.%

The amorphous properties were studied by electrical measurements. The sheet resistance was measured *in situ* using a four-point probe configuration, with the employment of a

devices.

for melt quenched amorphous films.

implanted film as due to compressive stress release.

atoms with distorted octahedral coordination.

**4. Doping by Ion beam: Amorphous properties**

during stress release.

and 6 at.% of fluorine.

thermal chuck.

Fig. 7. Conductivity versus reciprocal temperature for undoped amorphous GST (diamonds) and samples doped with oxygen (squares), nitrogen (triangles) or fluorine (circles).

Figure 7 shows the conductivity of amorphous samples as a function of the reciprocal temperature, measured in undoped GST (black diamonds) and in samples with: 1.5 at.% of O (full squares), 6 at.% of O (open squares), 1.5 at.% of N (full triangles) 10 at.% of N (open triangles), 1.5 at.% of F (full circles) and 6 at.% of F (open circles). As shown by the arrows in Fig.7, in the case of nitrogen or oxygen doping, the conductivity decreases as the dopant concentration increases. In the case of GST doped with fluorine, instead, the conductivity increases.

Fig. 8. Activation energy for conductivity as a function of doping concentration, obtained for nitrogen (diamonds), oxygen (squares) and fluorine (triangles).

Figure 8 shows the activation energy versus doping concentration, as obtained by linearly fitting data shown in Fig. 7.

In samples doped with nitrogen or oxygen the activation energy for conduction slightly increases, while for fluorine doping the activation energy remains almost the same. Largest variations have been obtained in the case of nitrogen: by doping GST with 10 at.% of N, the conductivity decreases of one order of magnitude and E increases from 0.42 *±* 0.01 eV, measured in undoped samples, up to 0.47 *±* 0.01 eV.

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 139

independent on time. This value indicates the establishment of steady state nucleation condition at the lowest grain radii. Moreover, the largest detected grain size increases with

Fig. 9. TEM image of undoped amorphous GST after anneling at 135°C for 50 min (a) and 66

From TEM images acquired at different annealing time, for several analyzed temperatures, it is possible to evaluate the evolution of the grain density ad the grain size distribution.

Fig. 10. Grain size distributions in undoped amorphous GST obtained after anneling at 130° C (a), 135°C (b) and 140°C (c). The distribution for each annealing temperature, have

where *r\* = 2 a* /*gv* is the critical radius, *a* is the lattice parameter, is the surface free energy at the interface, *Δgv* is the free energy difference between the two phases per atom, *f = f0 exp[Eg /(KT)]* is the thermally activated atomic jump frequency, and *Eg* is the activation

*gv /(KT) (1-r\*/r)* (3)

According to a well known approach based on the capillarity approximation, the growth or shrinkage of a cluster is due to the attachment or detachment of atoms at the interface between the two phases. Assuming the transition is limited by the reaction at the interface, the net growth rate of individual nuclei can be approximated to (Turnbull &. Fisher, 1949):

been obtained at different annealing time. Solid lines are fit to data.

*v(r) = a f*

time, following a linear law.

min (b).

These results are in agreement with changes in the optical band gaps of nitrogen doped GST films reported in literature (Kim et al., 2007), indicating that the increased band gap is closely related to the incorporation of nitrogen.

According to the results of numerical calculations shown by (Akola et al., 2010), an increased activation energy for conduction in nitrogen and oxygen doped samples could indicate the presence of more Ge atoms tetrahedrally coordinated.

## **5. Amorphous-to-crystal transition in undoped Ge2Sb2Te5 films**

Several experimental techniques have been employed to follow the amorphous-to-*fcc* transition in Ge2b2Te5, such as optical or electrical measurements. However, only microscopical techniques allow to separate the contribution of nucleation and growth, obtaining valuable information for understanding of the involved physical mechanisms. For this reason we have studied by *in situ* transmission electron microscopy (TEM) the amorphous to crystal phase transition as a function of temperature and annealing time, obtaining a detailed measurement of all the physical parameters governing the transformation.

To render the sample transparent for TEM, specimens were ground, dimpled and etched in a HNO3:HF:H2O (4:1:1) solution. This procedure has been adopted because it assures no variation of the initial interfaces and film thickness, which could affect the nucleation and growth process (Ohshima, 1996, Friedrich et al., 2000). Indeed, the specimen under observation in the microscope does not contain a hole, but it is made by a continuous film over its SiO2 substrate layer. Moreover, sample preparation by etching instead of conventional ion milling, can assures no undesired heating of the sample above the crystallization temperature, which could produce some pre-existing crystalline grains, therefore altering the kinetics of crystallization.

*In situ* TEM analyses have been performed using a JEOL 2010 TEM operating at 200 kV, with a sample holder equipped with a furnace around the specimen and a thermocouple. A magnification of 10000x was used for all the experiments. Since it has been shown the crystallization can be affected by the electron beam (Kooi et al., 2004), the beam has been switched off during annealing and fresh areas have been analysed, for each time step, in order to have a good statistical evaluation of the grain density and size. The transformation has been followed in the range of temperatures 130-143 °C.

Figure 9 (a) and (b) shows a TEM micrograph obtained for undoped GST after annealing at 135 °C for 50 min and 66 min, respectively. Grains appear as dark or bright circular regions and can be observed during the incubation regime (a), and after the onset of steady state nucleation (b). Grains with minimum detectable size of 7 nm are detected up to the complete crystallization, indicating that no saturation of nucleation sites occurs.

Figure 10 (a), (b), and (c) show the grain size distributions (GSDs) of samples annealed for various times, at 130 °C, 135°C and 140°C, respectively. GSDs have been obtained by assigning at each grain with surface area *A* an average radius *r = (A/)*, and by plotting, as a function of *r*, the density *dN* of crystalline grains with radius in the size interval *(r, r+dr)*, divided by *dr*. Figure 10 shows that all of the measured distributions smoothly decrease as a function of grain radius. At low grain size the distributions tend to reach the same value,

These results are in agreement with changes in the optical band gaps of nitrogen doped GST films reported in literature (Kim et al., 2007), indicating that the increased band gap is

According to the results of numerical calculations shown by (Akola et al., 2010), an increased activation energy for conduction in nitrogen and oxygen doped samples could

Several experimental techniques have been employed to follow the amorphous-to-*fcc* transition in Ge2b2Te5, such as optical or electrical measurements. However, only microscopical techniques allow to separate the contribution of nucleation and growth, obtaining valuable information for understanding of the involved physical mechanisms. For this reason we have studied by *in situ* transmission electron microscopy (TEM) the amorphous to crystal phase transition as a function of temperature and annealing time, obtaining a detailed measurement of all the physical parameters governing the

To render the sample transparent for TEM, specimens were ground, dimpled and etched in a HNO3:HF:H2O (4:1:1) solution. This procedure has been adopted because it assures no variation of the initial interfaces and film thickness, which could affect the nucleation and growth process (Ohshima, 1996, Friedrich et al., 2000). Indeed, the specimen under observation in the microscope does not contain a hole, but it is made by a continuous film over its SiO2 substrate layer. Moreover, sample preparation by etching instead of conventional ion milling, can assures no undesired heating of the sample above the crystallization temperature, which could produce some pre-existing crystalline grains,

*In situ* TEM analyses have been performed using a JEOL 2010 TEM operating at 200 kV, with a sample holder equipped with a furnace around the specimen and a thermocouple. A magnification of 10000x was used for all the experiments. Since it has been shown the crystallization can be affected by the electron beam (Kooi et al., 2004), the beam has been switched off during annealing and fresh areas have been analysed, for each time step, in order to have a good statistical evaluation of the grain density and size. The transformation

Figure 9 (a) and (b) shows a TEM micrograph obtained for undoped GST after annealing at 135 °C for 50 min and 66 min, respectively. Grains appear as dark or bright circular regions and can be observed during the incubation regime (a), and after the onset of steady state nucleation (b). Grains with minimum detectable size of 7 nm are detected up to the complete

Figure 10 (a), (b), and (c) show the grain size distributions (GSDs) of samples annealed for various times, at 130 °C, 135°C and 140°C, respectively. GSDs have been obtained by

a function of *r*, the density *dN* of crystalline grains with radius in the size interval *(r, r+dr)*, divided by *dr*. Figure 10 shows that all of the measured distributions smoothly decrease as a function of grain radius. At low grain size the distributions tend to reach the same value,

*(A/*

*)*, and by plotting, as

closely related to the incorporation of nitrogen.

therefore altering the kinetics of crystallization.

has been followed in the range of temperatures 130-143 °C.

crystallization, indicating that no saturation of nucleation sites occurs.

assigning at each grain with surface area *A* an average radius *r =*

transformation.

indicate the presence of more Ge atoms tetrahedrally coordinated.

**5. Amorphous-to-crystal transition in undoped Ge2Sb2Te5 films**

independent on time. This value indicates the establishment of steady state nucleation condition at the lowest grain radii. Moreover, the largest detected grain size increases with time, following a linear law.

Fig. 9. TEM image of undoped amorphous GST after anneling at 135°C for 50 min (a) and 66 min (b).

From TEM images acquired at different annealing time, for several analyzed temperatures, it is possible to evaluate the evolution of the grain density ad the grain size distribution.

Fig. 10. Grain size distributions in undoped amorphous GST obtained after anneling at 130° C (a), 135°C (b) and 140°C (c). The distribution for each annealing temperature, have been obtained at different annealing time. Solid lines are fit to data.

According to a well known approach based on the capillarity approximation, the growth or shrinkage of a cluster is due to the attachment or detachment of atoms at the interface between the two phases. Assuming the transition is limited by the reaction at the interface, the net growth rate of individual nuclei can be approximated to (Turnbull &. Fisher, 1949):

$$
\Delta v(r) = a \int \Delta \mathbf{g}\_v / (\mathbf{KT}) \tag{3}
$$

where *r\* = 2 a* /*gv* is the critical radius, *a* is the lattice parameter, is the surface free energy at the interface, *Δgv* is the free energy difference between the two phases per atom, *f = f0 exp[Eg /(KT)]* is the thermally activated atomic jump frequency, and *Eg* is the activation

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 141

calculated value from Eq (3) should be multiplied by an adimensional factor equal to 7 1013. For the steady-state nucleation rate *Jss* we have to consider an additional factor equal to 0.3 which, however, is reasonably explained by the evidence of heterogeneous nucleation. So, data can be well explained by employing the classical model of grain growth and nucleation based on capillarity, provided that we consistently assume an anomalously large

This result has been extensively explained in (Privitera et al., 2007a) by considering the effect of maximum entropy production rate, based on the large disorder of the *fcc* crystal structure. The role of entropy in the grain growth during first order phase transformations has been put in clear evidence in literature (Hill, 1990), for example driving the transition from planar to dendritic growth. It is reasonable to expect that in our case the entropy production rate is extremely large. In fact, it has been shown that the *fcc* structure of GST is a combination of two *fcc* lattices, one made of Te atoms and a second made of Ge, Sb, and vacancies, in the ratio 2:2:1 (Yamada et al., 1991). The Ge:Sb:vacancy *fcc* sublattice is randomly occupied by the three species. Kolobov et al. (Kolobov et al., 2004) have proposed a crystalline structure made by short-range ordered building blocks that may rotate by 90 degrees in an arbitrary direction, giving rise to a random distribution of Ge, Sb and vacancies on a long-range scale. Based on the type of bonding experimentally observed by EXAFS in amorphous and *fcc* phases, they have also proposed that the amorphous to *fcc* transition is likely due to an atomic flip of the Ge, which jumps from sp3 tetrahedral bonds with four Te atoms, to an octahedral configuration, again surrounded by Te first neighbors (eight) but with bonds of a prevalent p character. Kolobov proposes that the memory principle is based on this Ge flipping effect (from tetrahedral to octahedral). Since it involves at the ultimate level a single Ge atom and its immediate neighbors, this mechanism does possess a very large potential for memory scaling. However, this only effect, cannot explain the observed high transition rate, coupled with the high activation energy. What we have proposed is that the anomalous pre-exponential factors experimentally observed can be explained in terms of Ge flipping, but taking into account also the large configuration

Indeed, as a consequence of a Ge flip, it follows that the number of atoms crystallized in a single flip event is much larger than one. Referring to Fig. 5 in (Kolobov et al., 2004) this is at least equal to 6 Te atoms plus their first neighbors Ge, Sb, or vacancy. These amount to a total of about 18 atoms and vacancies crystallized for a single Ge flip. The Ge:Sb:vacancy *fcc* sublattice contains 11 positions, randomly occupied either by Sb atoms or by vacancies. Therefore it has a large number of different possible configurations. A rough estimate can be done considering that we have 11 allowed positions that can be occupied by either a vacancy or a Sb atom, each one with 4 different possible positions (Kolobov et al., 2004). This means that the total number of possible configurations is of the order of 811 8.6 109. So, according to (Hill, 1990) that has shown that the entropy generation rate is proportional to the growth velocity and that the growth mode in a phase transition follows the criterion of maximum entropy generation rate, the effective frequency of crystallization jumps should be multiplied by the factor 18 811 1.5 1011, taking into account for the number of

crystallized atoms per Ge flip ( 18) and for the configuration entropy term (811).

( 1014 times) factor multiplying the frequency of crystallization jumps.

entropy of the *fcc* crystal.

energy for the atomic motion responsible for the single elementary crystallization step. Therefore, for grain size much larger than *r\**, *v(r)* tends to *vg* = *a f gv /(KT).*

The nucleation process at any time *t* for a given cluster size *i* can be in general described considering the nucleation rate *Ji,t* of clusters with size larger than *i*.

It can be shown that the steady-state value of the nucleation rate (Becker & Döring, 1935) is:

$$I\_{ss} \equiv \mathcal{N}\_a \, O\_{l^\*} f \left[ \Delta \, g\_v / \left( 6 \,\pi \text{KT} \, ^\circ \text{I} \right) \right] 1^{\downarrow 2} \, \exp \left[ -\Delta \text{G} ^\circ / \left( \text{KT} \right) \right] \tag{4}$$

where *Na* is the atomic density and *Oi\* = 4 [3/(4)]2/3 i\* 2/3* is the number of atoms at the surface of a critical nucleus. Following this description, the steady-state nucleation rate *Jss* exhibits an activated temperature dependence with an activation energy *En*=*Eg* + *G\**.

The steady-state condition does not exist from the beginning of the transformation but it is established after a transient period during which the nucleation rate depends on size and time. All the data acquired from *in situ* TEM analyses as a function of annealing time and temperature, i.e. the grain size distributions, the transformed fraction and the grain density, have been simultaneously fitted following the model proposed by (Shneidman, 1987, 1988, 1991), as described in detail in (Privitera et al., 2007a). Four main parameters are used for the fitting procedure, the nucleation rate, the growth rate, the incubation time and the transient time , representing the time spent by the growing cluster in the critical region. Results of fitting have been shown in fig. 10 as solid lines.

The main parameters determined by fitting, i.e. the nucleation rate and the growth velocity, exhibit an activated temperature dependence, with activation energy of 2.6 eV and 2.3 eV, respectively (Privitera et al., 2007a).

In the classical theory of nucleation, the difference between the activation energy for the nucleation rate and that for the growth velocity is equal to the barrier energy for the nucleation *G\**, which corresponds to the energy necessary for the nucleation of a critical nucleus. Following this description, the value of *G\** here determined is less than 1eV and can be explained by considering the effect of heterogeneous nucleation. In this case, the model is also valid, provided that the atomic density *Na* is replaced by the density of nucleation sites and the barrier energy *G\** is reduced according to the lower surface energy term. Indeed, it has been observed that the nucleation of the metastable *fcc* phase is dependent on the substrate type and different activation energies have been found for Ge2Sb2Te5 over different substrates (Ohshima, 1996, Friedrich et al., 2000). To clarify this point in-situ isothermal annealing experiments of cross-sectional scanning-TEM (STEM) specimens of GST have been also performed (Lombardo et al. 2010). This analysis has shown the formation of crystalline grains preferentially occurs near the upper free interface, subsequently followed by the nucleation of crystal grain also at the GST / SiO2 interface.

By fitting the experimental data the pre-exponential terms can be also determined. Estimates of these can be done according to Eqs. (3) and (4) considering that *f0* is of the order of 1013 s-1, the lattice parameter *a* and the GST density *Na* are known, and the critical radius *r\** can be evaluated as the product of the measured *vg* at large sizes, multiplied by the transient time . Exponential factors for *Jss*, *vg* obtained from best fit to data, however, result anomalously different from the above estimates. In order to reproduce the experimental value of *vg,* the

energy for the atomic motion responsible for the single elementary crystallization step.

The nucleation process at any time *t* for a given cluster size *i* can be in general described

It can be shown that the steady-state value of the nucleation rate (Becker & Döring, 1935) is:

 *[3/(4*

exhibits an activated temperature dependence with an activation energy *En*=*Eg* +

surface of a critical nucleus. Following this description, the steady-state nucleation rate *Jss*

The steady-state condition does not exist from the beginning of the transformation but it is established after a transient period during which the nucleation rate depends on size and time. All the data acquired from *in situ* TEM analyses as a function of annealing time and temperature, i.e. the grain size distributions, the transformed fraction and the grain density, have been simultaneously fitted following the model proposed by (Shneidman, 1987, 1988, 1991), as described in detail in (Privitera et al., 2007a). Four main parameters are used for the fitting procedure, the nucleation rate, the growth rate, the incubation time and the transient time , representing the time spent by the growing cluster in the critical region.

The main parameters determined by fitting, i.e. the nucleation rate and the growth velocity, exhibit an activated temperature dependence, with activation energy of 2.6 eV and 2.3 eV,

In the classical theory of nucleation, the difference between the activation energy for the nucleation rate and that for the growth velocity is equal to the barrier energy for the

can be explained by considering the effect of heterogeneous nucleation. In this case, the model is also valid, provided that the atomic density *Na* is replaced by the density of

term. Indeed, it has been observed that the nucleation of the metastable *fcc* phase is dependent on the substrate type and different activation energies have been found for Ge2Sb2Te5 over different substrates (Ohshima, 1996, Friedrich et al., 2000). To clarify this point in-situ isothermal annealing experiments of cross-sectional scanning-TEM (STEM) specimens of GST have been also performed (Lombardo et al. 2010). This analysis has shown the formation of crystalline grains preferentially occurs near the upper free interface, subsequently followed by the nucleation of crystal grain also at the GST / SiO2 interface.

By fitting the experimental data the pre-exponential terms can be also determined. Estimates of these can be done according to Eqs. (3) and (4) considering that *f0* is of the order of 1013 s-1, the lattice parameter *a* and the GST density *Na* are known, and the critical radius *r\** can be evaluated as the product of the measured *vg* at large sizes, multiplied by the transient time . Exponential factors for *Jss*, *vg* obtained from best fit to data, however, result anomalously different from the above estimates. In order to reproduce the experimental value of *vg,* the

*G\**, which corresponds to the energy necessary for the nucleation of a critical

 *Jss=Na Oi\* f [Δ gv / (6 πKT i\* )] 1/2 exp[ -ΔG\* / (KT)]* (4)

*gv /(KT).*

*)]2/3 i\* 2/3* is the number of atoms at the

*G\** here determined is less than 1eV and

*G\** is reduced according to the lower surface energy

*G\**.

Therefore, for grain size much larger than *r\**, *v(r)* tends to *vg* = *a f*

where *Na* is the atomic density and *Oi\* = 4*

Results of fitting have been shown in fig. 10 as solid lines.

nucleus. Following this description, the value of

respectively (Privitera et al., 2007a).

nucleation sites and the barrier energy

nucleation

considering the nucleation rate *Ji,t* of clusters with size larger than *i*.

calculated value from Eq (3) should be multiplied by an adimensional factor equal to 7 1013. For the steady-state nucleation rate *Jss* we have to consider an additional factor equal to 0.3 which, however, is reasonably explained by the evidence of heterogeneous nucleation. So, data can be well explained by employing the classical model of grain growth and nucleation based on capillarity, provided that we consistently assume an anomalously large ( 1014 times) factor multiplying the frequency of crystallization jumps.

This result has been extensively explained in (Privitera et al., 2007a) by considering the effect of maximum entropy production rate, based on the large disorder of the *fcc* crystal structure. The role of entropy in the grain growth during first order phase transformations has been put in clear evidence in literature (Hill, 1990), for example driving the transition from planar to dendritic growth. It is reasonable to expect that in our case the entropy production rate is extremely large. In fact, it has been shown that the *fcc* structure of GST is a combination of two *fcc* lattices, one made of Te atoms and a second made of Ge, Sb, and vacancies, in the ratio 2:2:1 (Yamada et al., 1991). The Ge:Sb:vacancy *fcc* sublattice is randomly occupied by the three species. Kolobov et al. (Kolobov et al., 2004) have proposed a crystalline structure made by short-range ordered building blocks that may rotate by 90 degrees in an arbitrary direction, giving rise to a random distribution of Ge, Sb and vacancies on a long-range scale. Based on the type of bonding experimentally observed by EXAFS in amorphous and *fcc* phases, they have also proposed that the amorphous to *fcc* transition is likely due to an atomic flip of the Ge, which jumps from sp3 tetrahedral bonds with four Te atoms, to an octahedral configuration, again surrounded by Te first neighbors (eight) but with bonds of a prevalent p character. Kolobov proposes that the memory principle is based on this Ge flipping effect (from tetrahedral to octahedral). Since it involves at the ultimate level a single Ge atom and its immediate neighbors, this mechanism does possess a very large potential for memory scaling. However, this only effect, cannot explain the observed high transition rate, coupled with the high activation energy. What we have proposed is that the anomalous pre-exponential factors experimentally observed can be explained in terms of Ge flipping, but taking into account also the large configuration entropy of the *fcc* crystal.

Indeed, as a consequence of a Ge flip, it follows that the number of atoms crystallized in a single flip event is much larger than one. Referring to Fig. 5 in (Kolobov et al., 2004) this is at least equal to 6 Te atoms plus their first neighbors Ge, Sb, or vacancy. These amount to a total of about 18 atoms and vacancies crystallized for a single Ge flip. The Ge:Sb:vacancy *fcc* sublattice contains 11 positions, randomly occupied either by Sb atoms or by vacancies. Therefore it has a large number of different possible configurations. A rough estimate can be done considering that we have 11 allowed positions that can be occupied by either a vacancy or a Sb atom, each one with 4 different possible positions (Kolobov et al., 2004). This means that the total number of possible configurations is of the order of 811 8.6 109.

So, according to (Hill, 1990) that has shown that the entropy generation rate is proportional to the growth velocity and that the growth mode in a phase transition follows the criterion of maximum entropy generation rate, the effective frequency of crystallization jumps should be multiplied by the factor 18 811 1.5 1011, taking into account for the number of crystallized atoms per Ge flip ( 18) and for the configuration entropy term (811).

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 143

Samples doped with nitrogen exhibit higher resistance in the amorphous phase and the crystallization occurs at much higher temperatures than undoped material. A small increase in the crystallization temperature is also observed by oxygen doping while, on the contrary,

The dependence of the resistance as a function of time during isothermal anneals has been also studied by *in situ* resistance measurements using a four point probe. Figure 12 (a) and (b) shows the sheet resistance vs time as obtained during isothermal treatments at several annealing temperatures in the range 127-190°C, for undoped GST and for film doped by 6 at.% of fluorine, respectively. Fluorine doped samples are characterized by fast

Fig. 12. Sheet resistance versus time as measured during anneal at various temperatures in

By taking the first derivative of the resistance versus time, the characteristic time for transition can be measured, as the time at which a minimum in the derivative is observed. The characteristic time has been then measured for all the studied samples and it has been used to determine the temperature dependence of the transformation. The activation energy obtained for undoped Ge2Sb2Te5 is around 3 eV. The same value, within experimental errors, has been also determined for samples with low concentration of nitrogen and fluorine, as shown in Figure 13, reprinted from (Privitera et al. 2007b). Higher activation

the range 127-190°C, for undoped (a) and fluorine doped GST (b).

energy has been instead measured in GST doped with 6 at. % of oxygen.

addition of fluorine produces a decrease in the crystallization temperature.

crystallization.

This estimate does not reach the experimentally found factor (1.5 1011 against 7 1013) but it represents a lower limit, found assuming one single Ge flip, no effect of Sb, and that no other Ge neighbor is in the *fcc* crystal phase. If we consider the configuration entropy contribution due to the presence of the pre-existing amorphous / crystal interface, the estimate of the pre-exponential factor would be larger, so we believe that the estimate presented, strictly based on a single Ge flip, is satisfactorily close to the experimental result.

The discussion here presented can be considered a starting point, from which it is clearly necessary to further investigate and find other evidences in favor. In agreement with the proposed description, it has been observed that the crystallization speed is increased in GeTe-Sb2Te3 compounds by partially replacing Sb atoms with Bi atoms (Matsunaga & Yamada, 2004), or when a proportion of Ge is replaced by Sn (Kojima & Yamada, 2001). Since both Sn and Bi occupy the same (4b) position as Ge and Sb in the *fcc* lattice (Matsunaga & Yamada, 2004), their presence definitely increases the number of possible configurations and therefore the entropy factor in the growth velocity.

Recently, the disorder of phase change materials has been also taken into account to explain metal-semiconductor transitions occurring in crystalline phase change GeSbTe alloy (Siegrist et al., 2011). The strong disorder, usually associated only with amorphous solids, seems to be at the origin of the remarkable reproducibility of the resistance switching observed in phase change material and strongly suggests that the disorder could be one of the key ingredient that makes the physical properties of phase change materials so peculiar.

## **6. Amorphous to crystal transition in doped films**

Once we have studied in detail the crystallization process in undoped materials, we have investigated the effect of doping with light atom on the amorphous-to-crystal transition. In order to evidence the modifications due to doping, we have plotted in Fig. 11 the sheet resistance measured after 12 minutes anneals as a function of annealing temperature.

Fig. 11. Sheet resistance versus temperature as measured after 12 min isothermal anneals, for undoped and doped GST with oxygen (crosses), nitrogen (squares and triangles) or fluorine (circles).

This estimate does not reach the experimentally found factor (1.5 1011 against 7 1013) but it represents a lower limit, found assuming one single Ge flip, no effect of Sb, and that no other Ge neighbor is in the *fcc* crystal phase. If we consider the configuration entropy contribution due to the presence of the pre-existing amorphous / crystal interface, the estimate of the pre-exponential factor would be larger, so we believe that the estimate presented, strictly based on a single Ge flip, is satisfactorily close to the experimental result. The discussion here presented can be considered a starting point, from which it is clearly necessary to further investigate and find other evidences in favor. In agreement with the proposed description, it has been observed that the crystallization speed is increased in GeTe-Sb2Te3 compounds by partially replacing Sb atoms with Bi atoms (Matsunaga & Yamada, 2004), or when a proportion of Ge is replaced by Sn (Kojima & Yamada, 2001). Since both Sn and Bi occupy the same (4b) position as Ge and Sb in the *fcc* lattice (Matsunaga & Yamada, 2004), their presence definitely increases the number of possible configurations

Recently, the disorder of phase change materials has been also taken into account to explain metal-semiconductor transitions occurring in crystalline phase change GeSbTe alloy (Siegrist et al., 2011). The strong disorder, usually associated only with amorphous solids, seems to be at the origin of the remarkable reproducibility of the resistance switching observed in phase change material and strongly suggests that the disorder could be one of the key ingredient that makes the physical properties of phase change materials so peculiar.

Once we have studied in detail the crystallization process in undoped materials, we have investigated the effect of doping with light atom on the amorphous-to-crystal transition. In order to evidence the modifications due to doping, we have plotted in Fig. 11 the sheet

Fig. 11. Sheet resistance versus temperature as measured after 12 min isothermal anneals, for undoped and doped GST with oxygen (crosses), nitrogen (squares and triangles) or fluorine

resistance measured after 12 minutes anneals as a function of annealing temperature.

and therefore the entropy factor in the growth velocity.

**6. Amorphous to crystal transition in doped films**

(circles).

Samples doped with nitrogen exhibit higher resistance in the amorphous phase and the crystallization occurs at much higher temperatures than undoped material. A small increase in the crystallization temperature is also observed by oxygen doping while, on the contrary, addition of fluorine produces a decrease in the crystallization temperature.

The dependence of the resistance as a function of time during isothermal anneals has been also studied by *in situ* resistance measurements using a four point probe. Figure 12 (a) and (b) shows the sheet resistance vs time as obtained during isothermal treatments at several annealing temperatures in the range 127-190°C, for undoped GST and for film doped by 6 at.% of fluorine, respectively. Fluorine doped samples are characterized by fast crystallization.

Fig. 12. Sheet resistance versus time as measured during anneal at various temperatures in the range 127-190°C, for undoped (a) and fluorine doped GST (b).

By taking the first derivative of the resistance versus time, the characteristic time for transition can be measured, as the time at which a minimum in the derivative is observed. The characteristic time has been then measured for all the studied samples and it has been used to determine the temperature dependence of the transformation. The activation energy obtained for undoped Ge2Sb2Te5 is around 3 eV. The same value, within experimental errors, has been also determined for samples with low concentration of nitrogen and fluorine, as shown in Figure 13, reprinted from (Privitera et al. 2007b). Higher activation energy has been instead measured in GST doped with 6 at. % of oxygen.

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 145

Under annealing at higher temperature, samples doped with oxygen exhibit higher crystallization rate, compared to undoped GST. Figure 15 (a) and (b) shows TEM images of samples annealed in situ at 137°C for undoped and 6 at.% oxygen doped GST, respectively. Figure 15 (a) has been obtained in undoped material after 107 min annealing, while the image shown in fig. 15 (b) has been acquired on oxygen doped sample after only 70 minutes. Although annealed for shorter time, doped sample exhibits higher grain density and larger

Fig. 15. TEM image of undoped amorphous GST after anneling at 137°C for 107 min (a) and

Such a results is in agreement with the higher activation energy obtained in O doped blanket material (Fig. 13) and in oxygen doped memory cells (Matsuzaki et al., 2005). Moreover, increase of the crystal grain size has been also reported in 6 at. % oxygen-doped phase change optical recording disk, improving the overwrite cyclability of the optical recording disk. Therefore, the increase of activation energy, coupled with the increase of incubation time, makes oxygen doped GST particularly interesting for cell endurance improvement, without

The addition of 1.5 at.% of nitrogen by ion implantation produces strong variations in the

Figure 16 (a) and (b) shows the comparison between undoped sample (a) and sample with 1.5 at.% of N (b) after annealing at 140 °C for 75 min and 150 min, respectively. Nitrogen doped samples exhibit much longer incubation time compared to undoped film: after 150 min the material is still almost completely amorphous and only small crystallites are visible. Even after annealing at higher temperature, the morphology of nitrogen doped films appears different from the one observed in undoped materials. As an example, Figure 17 shows a TEM image of a nitrogen doped sample (1.5 at.%) annealed at 150°C for 80 minutes.

We have already observed in section 5 that the crystallization of undoped GST is heterogeneous and preferentially occurs near the upper free interface and, secondarily, at

of oxygen doped GST (6at.%) after annealing at 137°C for 70 min (b).

decreasing the crystallization rate, i.e. without increasing the writing time.

The film is crystalline but it is characterized by small crystalline size.

kinetics of the transformation.

the GST / SiO2 interface.

grain size.

Characteristic times observed in N and O doped samples are up to one order of magnitude higher than in pure GST, whilst the characteristic times observed in fluorine doped samples are lower.

Fig. 13. Activation energy for crystallization in Undoped and Doped GST with O, N, or F. Reprinted from (Privitera et al., 2007b)

Results from electrical measurements have been related to the kinetics of the transformation through *in situ* TEM analyses. Figure 14 shows typical TEM plan view images obtained during the amorphous-to-*fcc* transition in different samples annealed at 130°C. Fig. 14 (a) has been acquired for undoped sample annealed at 130°C for 240 min. A few crystalline grains embedded in the amorphous matrix are visible as dark or bright regions. Fig. 14 (b) shows a sample doped by 6 at.% of oxygen, annealed at the same temperature and time of 130°C 240 min. The material is almost completely amorphous and only small crystalline grains can be detected. Such a result is a confirmation of the observed increase of incubation and characteristic times observed by electrical measurements, as shown in figures 12 and 13.

Fig. 14. Comparison between TEM image of amorphous GST after anneling at 130°C for 240 min for undoped (a) and oxygen doped GST (6at.%) (b).

Characteristic times observed in N and O doped samples are up to one order of magnitude higher than in pure GST, whilst the characteristic times observed in fluorine doped samples

Fig. 13. Activation energy for crystallization in Undoped and Doped GST with O, N, or F.

Results from electrical measurements have been related to the kinetics of the transformation through *in situ* TEM analyses. Figure 14 shows typical TEM plan view images obtained during the amorphous-to-*fcc* transition in different samples annealed at 130°C. Fig. 14 (a) has been acquired for undoped sample annealed at 130°C for 240 min. A few crystalline grains embedded in the amorphous matrix are visible as dark or bright regions. Fig. 14 (b) shows a sample doped by 6 at.% of oxygen, annealed at the same temperature and time of 130°C 240 min. The material is almost completely amorphous and only small crystalline grains can be detected. Such a result is a confirmation of the observed increase of incubation and characteristic times observed by electrical measurements, as shown in figures 12 and 13.

Fig. 14. Comparison between TEM image of amorphous GST after anneling at 130°C for 240

min for undoped (a) and oxygen doped GST (6at.%) (b).

are lower.

Reprinted from (Privitera et al., 2007b)

Under annealing at higher temperature, samples doped with oxygen exhibit higher crystallization rate, compared to undoped GST. Figure 15 (a) and (b) shows TEM images of samples annealed in situ at 137°C for undoped and 6 at.% oxygen doped GST, respectively. Figure 15 (a) has been obtained in undoped material after 107 min annealing, while the image shown in fig. 15 (b) has been acquired on oxygen doped sample after only 70 minutes. Although annealed for shorter time, doped sample exhibits higher grain density and larger grain size.

Fig. 15. TEM image of undoped amorphous GST after anneling at 137°C for 107 min (a) and of oxygen doped GST (6at.%) after annealing at 137°C for 70 min (b).

Such a results is in agreement with the higher activation energy obtained in O doped blanket material (Fig. 13) and in oxygen doped memory cells (Matsuzaki et al., 2005). Moreover, increase of the crystal grain size has been also reported in 6 at. % oxygen-doped phase change optical recording disk, improving the overwrite cyclability of the optical recording disk. Therefore, the increase of activation energy, coupled with the increase of incubation time, makes oxygen doped GST particularly interesting for cell endurance improvement, without decreasing the crystallization rate, i.e. without increasing the writing time.

The addition of 1.5 at.% of nitrogen by ion implantation produces strong variations in the kinetics of the transformation.

Figure 16 (a) and (b) shows the comparison between undoped sample (a) and sample with 1.5 at.% of N (b) after annealing at 140 °C for 75 min and 150 min, respectively. Nitrogen doped samples exhibit much longer incubation time compared to undoped film: after 150 min the material is still almost completely amorphous and only small crystallites are visible.

Even after annealing at higher temperature, the morphology of nitrogen doped films appears different from the one observed in undoped materials. As an example, Figure 17 shows a TEM image of a nitrogen doped sample (1.5 at.%) annealed at 150°C for 80 minutes. The film is crystalline but it is characterized by small crystalline size.

We have already observed in section 5 that the crystallization of undoped GST is heterogeneous and preferentially occurs near the upper free interface and, secondarily, at the GST / SiO2 interface.

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 147

Fig. 18. Bright field TEM micrographs in cross section of a sample doped by N 1.5 at.% asdeposited (a), and annealed in situ at 150°C for 20 min (b). The first crystalline formation is clearly visible at the GST/SiO2 interface, with additional crystalline grains nucleated also at

Therefore it appears clear from the experimental data that, at least in the range of compositions here investigated, the oxygen or nitrogen doping affects the crystallization process differently, the first one producing higher crystallization rate and larger grain size,

The increase of the crystallization temperature in oxygen (Jeong et al., 2001) and nitrogen [Kojima et al., 1998, Jeong et al., 2000) doped GST has been already reported in literature, as well as the reduction of the grain size in samples containing nitrogen. Several models have been proposed (Jeong et al., 2001, Jeong et al., 2000) trying to simultaneously explaining the effect of both O and N in order to find the optimal dopant to engineer

First, it has been observed that the reduced grain size could be caused by dopant segregation near grain boundaries in the form of nitrides (Kojima et al., 1998). The resistance of the crystalline phase would increase due to additional electron scattering from grain boundaries (Horii et al., 2003). We cannot exclude, in the samples with 10 at.% of N, the formation of very thin nitrides layers at grain boundaries, which are probably responsible for the increased resistivity of the crystalline phase and for the smoothing of the shape of the

Regarding oxygen doping, the formation of Ge or Sb oxides has been reported for concentrations above 10 at.% (Jang et al., 2009, Morales-Sanchez et al., 2006). For lower concentration, as in the samples we have studied, no oxidation seems to occur. It has been therefore proposed that oxygen atoms, as well as nitrogen atoms at low concentration, are located at the tetrahedral interstitial sites (Jeong et al., 2000). Indeed, since nitrogen and oxygen have quite small atomic sizes compared to that of Ge, Sb and Te, they cannot substitute these atoms, filling the vacancies which are present in the structure. Therefore their effect is expected to be very different from the one observed in the case of Bi or Sn, (Matsunaga & Yamada, 2004), which substituting in the Ge:Sb:vacancy lattice, can increase the number of possible configurations and therefore the growth velocity. N or O atoms will be instead located at the tetrahedral interstitial sites distorting the unit cell and resulting in a strain field (Jeong et al., 2000). Since nitrogen has a larger atomic radius (0.071 nm) (Lide, 2004), compared to oxygen (0.066 nm) (Lide, 2004), a larger amount of strain is expected to

Another factor to be taken into account together with the stress is the chemical bond between Ge and Te and its modifications due to the presence of nitrogen, or other dopants.

while N doping gives rise to grain size refinement and longer incubation time.

the upper interface.

specific properties of GST.

resistance versus temperature shown in Fig. 11.

accommodate nitrogen atoms in the structure.

Fig. 16. TEM image of undoped amorphous GST after anneling at 140°C for 75 min (a) and of nitrogen doped GST (1.5 at.%) after annealing at 140°C for 150 min (b).

Since the ion implantation can modify the interfaces, causing some intermixing, we have performed in situ TEM analyses in cross section. Figure 18 (a) and (b) shows the amorphous GST film doped with 1.5 at.% of N as deposited and after annealing in situ at 150°C for 20 min, respectively. As observed for undoped GST (Lombardo et al., 2010) a continuous polycrystalline layer is formed. However, in contrast with observation on pure GST, in the case of nitrogen doping the nucleation has occurred preferentially at the GST/SiO2 interface, with the addition of small dispersed crystalline regions in the amorphous material and at the upper interface. Therefore doping by ion implantation is also effective in modifying the nucleation sites at which heterogeneous nucleation preferentially occurs.

Fig. 17. TEM image of a sample doped by 1.5at.%of nitrogen annealed at 150°C for 80 min.

Fig. 16. TEM image of undoped amorphous GST after anneling at 140°C for 75 min (a) and

Since the ion implantation can modify the interfaces, causing some intermixing, we have performed in situ TEM analyses in cross section. Figure 18 (a) and (b) shows the amorphous GST film doped with 1.5 at.% of N as deposited and after annealing in situ at 150°C for 20 min, respectively. As observed for undoped GST (Lombardo et al., 2010) a continuous polycrystalline layer is formed. However, in contrast with observation on pure GST, in the case of nitrogen doping the nucleation has occurred preferentially at the GST/SiO2 interface, with the addition of small dispersed crystalline regions in the amorphous material and at the upper interface. Therefore doping by ion implantation is also effective in modifying the

Fig. 17. TEM image of a sample doped by 1.5at.%of nitrogen annealed at 150°C for 80 min.

of nitrogen doped GST (1.5 at.%) after annealing at 140°C for 150 min (b).

nucleation sites at which heterogeneous nucleation preferentially occurs.

Fig. 18. Bright field TEM micrographs in cross section of a sample doped by N 1.5 at.% asdeposited (a), and annealed in situ at 150°C for 20 min (b). The first crystalline formation is clearly visible at the GST/SiO2 interface, with additional crystalline grains nucleated also at the upper interface.

Therefore it appears clear from the experimental data that, at least in the range of compositions here investigated, the oxygen or nitrogen doping affects the crystallization process differently, the first one producing higher crystallization rate and larger grain size, while N doping gives rise to grain size refinement and longer incubation time.

The increase of the crystallization temperature in oxygen (Jeong et al., 2001) and nitrogen [Kojima et al., 1998, Jeong et al., 2000) doped GST has been already reported in literature, as well as the reduction of the grain size in samples containing nitrogen. Several models have been proposed (Jeong et al., 2001, Jeong et al., 2000) trying to simultaneously explaining the effect of both O and N in order to find the optimal dopant to engineer specific properties of GST.

First, it has been observed that the reduced grain size could be caused by dopant segregation near grain boundaries in the form of nitrides (Kojima et al., 1998). The resistance of the crystalline phase would increase due to additional electron scattering from grain boundaries (Horii et al., 2003). We cannot exclude, in the samples with 10 at.% of N, the formation of very thin nitrides layers at grain boundaries, which are probably responsible for the increased resistivity of the crystalline phase and for the smoothing of the shape of the resistance versus temperature shown in Fig. 11.

Regarding oxygen doping, the formation of Ge or Sb oxides has been reported for concentrations above 10 at.% (Jang et al., 2009, Morales-Sanchez et al., 2006). For lower concentration, as in the samples we have studied, no oxidation seems to occur. It has been therefore proposed that oxygen atoms, as well as nitrogen atoms at low concentration, are located at the tetrahedral interstitial sites (Jeong et al., 2000). Indeed, since nitrogen and oxygen have quite small atomic sizes compared to that of Ge, Sb and Te, they cannot substitute these atoms, filling the vacancies which are present in the structure. Therefore their effect is expected to be very different from the one observed in the case of Bi or Sn, (Matsunaga & Yamada, 2004), which substituting in the Ge:Sb:vacancy lattice, can increase the number of possible configurations and therefore the growth velocity. N or O atoms will be instead located at the tetrahedral interstitial sites distorting the unit cell and resulting in a strain field (Jeong et al., 2000). Since nitrogen has a larger atomic radius (0.071 nm) (Lide, 2004), compared to oxygen (0.066 nm) (Lide, 2004), a larger amount of strain is expected to accommodate nitrogen atoms in the structure.

Another factor to be taken into account together with the stress is the chemical bond between Ge and Te and its modifications due to the presence of nitrogen, or other dopants.

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 149

Ion implantation can be also adopted for doping of phase change materials, modifying the amorphous-to-crystal transition kinetics, upon which the reliability definitely depends. The crystallization is a fast process occurring at very low temperature, compared to other solid materials. The explanation of such a fast transition has been found in the peculiar structure of the *fcc* phase which contains a large amount of disorder and whose formation is driven by flips of Ge atoms. By nitrogen implantation it is possible to increase the crystallization temperature and the crystalline resistance improving the cell endurance, reducing the RESET current and enabling the use of PCM also for high temperature applications, such as automotive. The formation of Ge-N bonds has been identified as the main effect determining the increase of crystallization temperature, the increase of the band gap in the

By oxygen implantation, a moderate increase of crystallization temperature has been obtained, maintaining short crystallization time, thanks to the increased activation energy. Decrease of crystallization temperature has been instead obtained by fluorine doping. The way these doping acts is still not well understood and more extensive studies are required to relate the chemical nature of dopants belonging to different groups in the periodic table to the induced modifications in the phase change structure, in order to optimize the dopant

The author wish to gratefully acknowledge Prof. Emanuele Rimini (University of Catania), Corrado Bongiorno (CNR), Romina Zonca (STMicroelectronics), Agostino Pirovano

Akola, J.; Larrucea, J. & Jones, R. O. (2010), Polymorphism in amorphous Ge2Sb2Te5 :

B.-S. Lee, G. W. Burr, R. M. Shelby, S. Raoux, C. T. Rettner, S. N. Bogle, K. Darmawikarta, S.

Baeri, P. & Campisano, S.U. (1982) Chapter 4: Heat Flow Calculations, In: *Laser Annealing of*

Becker, R. & Döring, W. (1935) Kinetische behandlung der keimbildung in übersättigten

Boniardi, M.; Redaelli, A.; Pirovano, A.; Tortorelli, I.; Ielmini, D.; & Pellizzer, F. (2009). A

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*Semiconductors,* ed. by J. M. Poate & J. W. Mayer, pp. 75-110, Academic Press, ISBN

dämpfen, *Annalen der Physik* (Leipzig), Vol. 24, N. 8 (1935), pp.719-752, ISSN

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amorphous material and the reduction of crystalline grain size.

(Numonyx) and Roberto Bez (Numonyx) for the fruitful collaboration.

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Italy, 2010. Available from

980-984, ISSN 0036-8075.

0125588208, New York.

084506-5, ISSN 0021-8979.

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**8. Acknowledgment**

**9. References**

The main crystallization mechanism has been identified as an atomic flip of the Ge, which jumps from tetrahedral bonds with 4 Te atoms, to an octahedral configuration, surrounded by 8 Te first neighbors. Xray photoemission spectroscopy (XPS) analyses by employing synchrotron radiation have shown that when a GST film is crystallized in the *fcc* structure, the intensity of Germanium in tetrahedral position is decreased while the intensity of Ge octahedral is increased owing to the umbrella-flip change in Ge atoms from the tetrahedral to the octahedral position (Kolobov et al., 2004).

Both experimental data (Kim et al., 2007, Jung et al., 2007) and atomic *ab initio* calculations (Cho et al., 2011) performed on nitrogen doped GST have shown that the N atoms are prevalently bonded with the Ge atoms to form GeNx, rather than bonding with the Te or Sb atoms. Moreover, data of Ref. (Kim et al., 2007) show that the formation of Ge-N bonds can disturb the umbrella-flip of Ge atoms between the tetrahedral and octahedral positions because, as nitrogen is incorporated into the GST film, the intensity of the peak relative to Germanium in the octahedral position is decreased, since nitrogen combines with Ge to form germanium nitrides. By reducing the portion of the peak of Ge octahedral nitrogen is therefore effective in suppressing the crystallization.

A similar behavior is expected also for oxygen atoms, with the formation of Ge-O bonds. However, nitrogen can form three bonds while oxygen only two. Moreover, N dopants increase the portion of tetrahedral Ge atoms, while, on the base of calculations reported in (Cho et al., 2011), O doped GST seems to be more similar to undoped material, without favoring tetrahedral sites respect to octahedral Ge atoms. This could explain the lower increase in the energy gap observed in the case of oxygen (Fig. 8).

Considering the heterogeneous character of the transformation, the proposed assumption that oxygen and a small amount of germanium oxide could act as nucleation centers during crystallization, increasing the crystallization rate (Morales-Sanchez et al., 2006) could be a good explanation.

The decrease of crystallization temperature observed for fluorine doping remains to be better investigated. It has been reported that amorphous GST can have fast crystallization under compressive stress. So the effect of F could be explained with the introduction of compressive stress in the amorphous material. More extensive studies however are needed to map out the relationship between the chemical nature of dopants and the induced modifications in the phase change structure, in order to better identify the optimal dopant in order to achieve the desired properties.

## **7. Conclusions**

In this chapter we reviewed some applications of ion beam to the study of the physical properties of Ge2Sb2Te5, the phase change material most widely employed for application as nonvolatile memories. Phase change memories are one of the most promising solutions as replacement of current FLASH technology, based on charge storage, but some material properties need to be optimized. We have shown that ion implantation can be successfully employed to modify the properties of the amorphous phase change material, obtaining conduction properties very similar to those of melt-quenched amorphous material, formed by electric pulses or laser in memory cells.

Ion implantation can be also adopted for doping of phase change materials, modifying the amorphous-to-crystal transition kinetics, upon which the reliability definitely depends. The crystallization is a fast process occurring at very low temperature, compared to other solid materials. The explanation of such a fast transition has been found in the peculiar structure of the *fcc* phase which contains a large amount of disorder and whose formation is driven by flips of Ge atoms. By nitrogen implantation it is possible to increase the crystallization temperature and the crystalline resistance improving the cell endurance, reducing the RESET current and enabling the use of PCM also for high temperature applications, such as automotive. The formation of Ge-N bonds has been identified as the main effect determining the increase of crystallization temperature, the increase of the band gap in the amorphous material and the reduction of crystalline grain size.

By oxygen implantation, a moderate increase of crystallization temperature has been obtained, maintaining short crystallization time, thanks to the increased activation energy. Decrease of crystallization temperature has been instead obtained by fluorine doping. The way these doping acts is still not well understood and more extensive studies are required to relate the chemical nature of dopants belonging to different groups in the periodic table to the induced modifications in the phase change structure, in order to optimize the dopant specie on the base of the desired properties.

## **8. Acknowledgment**

The author wish to gratefully acknowledge Prof. Emanuele Rimini (University of Catania), Corrado Bongiorno (CNR), Romina Zonca (STMicroelectronics), Agostino Pirovano (Numonyx) and Roberto Bez (Numonyx) for the fruitful collaboration.

## **9. References**

148 Ion Implantation

The main crystallization mechanism has been identified as an atomic flip of the Ge, which jumps from tetrahedral bonds with 4 Te atoms, to an octahedral configuration, surrounded by 8 Te first neighbors. Xray photoemission spectroscopy (XPS) analyses by employing synchrotron radiation have shown that when a GST film is crystallized in the *fcc* structure, the intensity of Germanium in tetrahedral position is decreased while the intensity of Ge octahedral is increased owing to the umbrella-flip change in Ge atoms from the tetrahedral

Both experimental data (Kim et al., 2007, Jung et al., 2007) and atomic *ab initio* calculations (Cho et al., 2011) performed on nitrogen doped GST have shown that the N atoms are prevalently bonded with the Ge atoms to form GeNx, rather than bonding with the Te or Sb atoms. Moreover, data of Ref. (Kim et al., 2007) show that the formation of Ge-N bonds can disturb the umbrella-flip of Ge atoms between the tetrahedral and octahedral positions because, as nitrogen is incorporated into the GST film, the intensity of the peak relative to Germanium in the octahedral position is decreased, since nitrogen combines with Ge to form germanium nitrides. By reducing the portion of the peak of Ge octahedral nitrogen is

A similar behavior is expected also for oxygen atoms, with the formation of Ge-O bonds. However, nitrogen can form three bonds while oxygen only two. Moreover, N dopants increase the portion of tetrahedral Ge atoms, while, on the base of calculations reported in (Cho et al., 2011), O doped GST seems to be more similar to undoped material, without favoring tetrahedral sites respect to octahedral Ge atoms. This could explain the lower

Considering the heterogeneous character of the transformation, the proposed assumption that oxygen and a small amount of germanium oxide could act as nucleation centers during crystallization, increasing the crystallization rate (Morales-Sanchez et al., 2006) could be a

The decrease of crystallization temperature observed for fluorine doping remains to be better investigated. It has been reported that amorphous GST can have fast crystallization under compressive stress. So the effect of F could be explained with the introduction of compressive stress in the amorphous material. More extensive studies however are needed to map out the relationship between the chemical nature of dopants and the induced modifications in the phase change structure, in order to better identify the optimal dopant in

In this chapter we reviewed some applications of ion beam to the study of the physical properties of Ge2Sb2Te5, the phase change material most widely employed for application as nonvolatile memories. Phase change memories are one of the most promising solutions as replacement of current FLASH technology, based on charge storage, but some material properties need to be optimized. We have shown that ion implantation can be successfully employed to modify the properties of the amorphous phase change material, obtaining conduction properties very similar to those of melt-quenched amorphous material, formed

to the octahedral position (Kolobov et al., 2004).

therefore effective in suppressing the crystallization.

good explanation.

**7. Conclusions**

order to achieve the desired properties.

by electric pulses or laser in memory cells.

increase in the energy gap observed in the case of oxygen (Fig. 8).

Akola, J.; Larrucea, J. & Jones, R. O. (2010), Polymorphism in amorphous Ge2Sb2Te5 : comparison of melt-quenched and as-deposited structures, *EPCOS 2010*, Milano, Italy, 2010. Available from

http://www.epcos.org/library/papers/pdf\_2010/Oral/F03-Akola.pdf

Ion Implantation in Phase Change Ge2Sb2Te5 Thin Films for Non Volatile Memory Applications 151

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**8**

Irina Antonova

*Russia*

**Si Nanocrystal Arrays Created in SiO2**

*A.V. Rzhanov Institute of Semiconductor Physics SB RAS*

**Matrix by High-Energy Ion Bombardment**

In the next two decades, nanoscience and nanotechnologies will transfer into new products and processes. One emerging area where this challenge will be successfully met is the field of semiconductor nanocrystals (NCs) (see, for instance, Norris et al., 2008). The promises offered by dielectric layers with hosted semiconductor nanocrystals in applications include production of silicon-based light emitting systems (Iacona et al., 2006), wavelength-tunable lasers (Klimov et al., 2007), memory elements (Hu et al., 2009), solar cells (Gur et al., 2008), low-dimensional devices (for instance, single electron transistors operating at room temperature (Zhou et al., 2010)), and a wide spectrum of biotechnologies (for instance, bioimaging (Michalet et al., 2005)). The prospects for developing new electronic and optoelectronic devices on the basis of spatially ordered arrays of quantum dots (QDs) embedded in a semiconductor or insulator matrix have arouse currently growing interest in the production methods of such arrays and in the

The idea to use non-overlapping tracks of high-energy (*E* ≥ 1 MeV/amu) heavy ions for production of nanocrystalline (NC) systems is becoming increasingly attractive among the various methods for fabrication of nanodimensional structures (nanolithography, molecular nanoelectronics, organic nanostructures, etc.) (Fink et al., 2007; Rizza et al., 2009). Beams of heavy fast ions can be efficiently used not only for formation, but also for modification, of nanodimensional structures, in particular, metallic nanoclusters (Klaumunzer, 2006; D'Orleans et al., 2003). Avasthi et al. (Avasthi et al., 2010) have summarized experimental data on the formation of elongated metal nanoparticles, and suggested a model for the interaction of fast ions with such nanoparticles dependently on nanoparticle sizes. Studies of the effect of high-energy ion irradiation on the formation of semiconductor (silicon or germanium) nanoclusters have also appeared recently (Chaudhari et al., 2003; 2005; Arnoldbik et al., 2005). It was shown that, as a consequence of the high density of ionization energy losses in the material, silicon nanoinclusions could be formed in SiO*<sup>x</sup>* layers without any additional thermal treatment given to the samples. Studies of the effect of high-energy ion irradiation on SiO2 layers containing Si nanocrystals or an excess concentration of dissolved silicon (Si:SiO2 layers) have made it possible to propose a new approach to formation of vertically ordered NC arrays (Antonova et al., 2008 a; 2009 a;

**1. Introduction**

2009 b; 2010 a; 2010 b).

studies of their optical and electronic properties.

## **Si Nanocrystal Arrays Created in SiO2 Matrix by High-Energy Ion Bombardment**

Irina Antonova

*A.V. Rzhanov Institute of Semiconductor Physics SB RAS Russia*

## **1. Introduction**

152 Ion Implantation

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In the next two decades, nanoscience and nanotechnologies will transfer into new products and processes. One emerging area where this challenge will be successfully met is the field of semiconductor nanocrystals (NCs) (see, for instance, Norris et al., 2008). The promises offered by dielectric layers with hosted semiconductor nanocrystals in applications include production of silicon-based light emitting systems (Iacona et al., 2006), wavelength-tunable lasers (Klimov et al., 2007), memory elements (Hu et al., 2009), solar cells (Gur et al., 2008), low-dimensional devices (for instance, single electron transistors operating at room temperature (Zhou et al., 2010)), and a wide spectrum of biotechnologies (for instance, bioimaging (Michalet et al., 2005)). The prospects for developing new electronic and optoelectronic devices on the basis of spatially ordered arrays of quantum dots (QDs) embedded in a semiconductor or insulator matrix have arouse currently growing interest in the production methods of such arrays and in the studies of their optical and electronic properties.

The idea to use non-overlapping tracks of high-energy (*E* ≥ 1 MeV/amu) heavy ions for production of nanocrystalline (NC) systems is becoming increasingly attractive among the various methods for fabrication of nanodimensional structures (nanolithography, molecular nanoelectronics, organic nanostructures, etc.) (Fink et al., 2007; Rizza et al., 2009). Beams of heavy fast ions can be efficiently used not only for formation, but also for modification, of nanodimensional structures, in particular, metallic nanoclusters (Klaumunzer, 2006; D'Orleans et al., 2003). Avasthi et al. (Avasthi et al., 2010) have summarized experimental data on the formation of elongated metal nanoparticles, and suggested a model for the interaction of fast ions with such nanoparticles dependently on nanoparticle sizes. Studies of the effect of high-energy ion irradiation on the formation of semiconductor (silicon or germanium) nanoclusters have also appeared recently (Chaudhari et al., 2003; 2005; Arnoldbik et al., 2005). It was shown that, as a consequence of the high density of ionization energy losses in the material, silicon nanoinclusions could be formed in SiO*<sup>x</sup>* layers without any additional thermal treatment given to the samples. Studies of the effect of high-energy ion irradiation on SiO2 layers containing Si nanocrystals or an excess concentration of dissolved silicon (Si:SiO2 layers) have made it possible to propose a new approach to formation of vertically ordered NC arrays (Antonova et al., 2008 a; 2009 a; 2009 b; 2010 a; 2010 b).

Si Nanocrystal Arrays Created in SiO2 Matrix by High-Energy Ion Bombardment 155

Research, Dubna. For the ion energies used, the ion penetration range was 15–23 μm (see Table 1); i.e. it was substantially larger than the thickness of silicon-rich oxide layers. This means that almost all ions penetrated into the substrate, and the defects generated in elastic scattering channels were also predominantly introduced in the substrate, whereas

With reference to SRIM data, the electron stopping losses due to ionization dEe/dx in the oxide layers were expected to vary from 7.1 (Kr), 10.2 (Xe), and 22.0 keV/nm (Bi) for the SiO2 end of the sample to 9.5 (Kr), 12.3 (Xe), and 24.3 keV/nm (Bi) for its Si end. The ionization losses of ion energy lead to ionization of atoms in the material under irradiation. In a time interval of 10-13 - 10-12 s the excitation energy was transferred into anisotropic heating of the oxide layers due to thermalization process. Using the value of ionization losses, the dependence of ion-track sizes in -SiO2 on ion energy, and the thermal spike model of (Toulemonde et al., 2000; Meftah et al., 1997), we could estimate the width of ion tracks in our samples, which turned out to be about 4-6 nm. Results for the time evolution of temperature inside ion tracks calculated by TSRIKE-02 software (Toulemonde et al., 2000) are shown in Fig.1a. TSPIKE-02 simulates the electron-lattice interaction (transfer of the ionization losses of ions into the heating of the target lattice). The variable parameter of the ion-lattice interaction is the radius of the shells surrounding individual ion trajectories in which the main portion of ion energy is lost for ionization. As it follows from Fig.1a (see also Fig.15 a), the local transient temperature in ion track cores in our experiments could reach 1700 - 2500oC, the ion-track heating time being about 10-11 - 10-10 s. It means that almost all inter-atomic bonds along ion tracks could be broken to free the atoms and make them capable of displacing from their regular lattice sites. As a result, new bonds could readily form in the material, this process leading to subsequent amorphization or crystallization of part of previously existing Si nanoparticles as well as to the formation of new NCs in the

One more effect that was taken into consideration for materials subjected to the irradiation with high-energy ions (see also Section 6) is ion-beam-induced anisotropic plastic deformations (Van Dillen et al., 2004; 2005). The strain anisotropy is imposed by the direction of the ion beam: normally, the material expands in the direction normal to the ion beam and contracts parallel to it, while maintaining their volume. According to the developed model, the strain (determined as the relative change of lattice constant) versus time goes through a maximum and then reaches a constant value. The anisotropic strain rises with ion energy and fluence. It should be emphasized here that up to fluences 1012 cm-2 the ion tracks in the material were independent of each other, whereas at higher fluences

The ionization losses of the ions and the anisotropic heating of the material are higher in Si in comparison with SiO2. The anisotropic deformations found in amorphous silicon are typically 1.5 orders of magnitude lower than the anisotropic deformations in SiO2 (Van Dillen et al., 2005; Meftah et al., 1997). This means that the anisotropic strain and the anisotropic heating should demonstrate considerable changes with increasing concentration of Si nanocrystals embedded in the SiO2 layer. In the case of variable NC concentration we have to expect a strong anisotropic-strain effect for SiO2 layers with low Si-phase content and a more pronounced anisotropic heating effect for SiO2 layers with

prevailing energy losses in the layers were the electron stopping losses.

oxide layer along ion tracks.

they can become overlapping.

high excess Si phase content.

In the present review, we demonstrate the possibility of using irradiation of samples with high-energy gas ions (Kr, Xe, and Bi) for introduction and modification of Si nanocrystals in Si:SiO2 and NC:SiO2 layers, and also for formation of ordered arrays of Si nanocrystals embedded in a SiO2 matrix. Due to the small size of nanocrystals synthesized in our samples (typically, ≤ 6 - 8 nm), in describing our results we will also use an alternative term 'quantum dots'. We will analyze the changes in the optical and electronic properties of NC systems and discuss possible applications of such systems. It should be emphasized here that a special feature of used irradiation, in contrast to conventional ion implantation, is huge ion-energy losses in the bulk of Si:SiO2 layers and simultaneous creation of a relatively low defect concentration in such layers (Antonova et al., 2006a; Toulemonde et al., 2000; Meftah et al., 1997). The factor causing a low defect concentration in examined irradiated layers is the small thickness of those layers, which is normally much smaller than the projected ion range. Hence, while in practice defects are introduced in the substrate of film systems, the energy is mainly dissipated in the top Si:SiO2 layer. The latter circumstance has allowed us to separate out the positive effect of ion irradiation (the huge energy losses in ion tracks) from the negative effect (defect formation).

## **2. Sample preparation and ion irradiation conditions**

#### **2.1 Nanocrystals in SiO2 layers with varied concentration of excess Si**

The Si:SiO2 layers of thickness 400–1000 nm used in our study were sputtered from Si and SiO2 sources onto silicon substrates. In the experimental arrangement used, the SiO2 and Si sources were spaced ~100 mm apart. The large distance between the sources provided a possibility to form layers with varied concentration of excess silicon in the oxide. After annealing of samples and formation of nanocrystals, we could obtain layers with concentration and sizes of NCs varied in the lateral plane. That gave us an interesting opportunity to examine the properties of formed Si:SiO2 layers as dependent on NC concentration and sizes. The excess Si content of the silicon-rich oxide layers ranged from 4 to 94 vol.%. The composition of deposited layers was estimated from the thickness of Si and SiO2 layers obtained by individual deposition of films from the Si and SiO2 sources onto different substrates (Shamin et al., 2010). Subsequent high-temperature annealing (at 1140°C for 40 min) used for part of samples resulted in the formation of silicon oxide layers containing Si nanocrystals (NC:SiO2 layers), whose sizes only slightly depended on layer composition and whose concentration increased with coordinate along the sample by several orders of magnitude (Sa'ar et al., 2005; Antonova et al., 2006b). Experimental results are presented below as a function of excess Si concentration. This parameter is used to characterize the composition of our Si:SiO2 layers since its value was inherently defined by the evaporation procedure, whereas other parameters, such as NC concentration and sizes, were complicated functions of layer composition, and also of thermal treatments given to the samples.

#### **2.2 High-energy ion irradiation: Specific features**

For irradiation of Si:SiO2 and NC:SiO2 layers obtained as described above, Kr, Xe, or Bi ions respectively with energies 90, 130-250, and 65 - 670 MeV were used. The ion fluences ranged from 3x1011 to 3x1013 cm–2. The irradiations of our samples were performed with the use of accelerators available in the Laboratory of Nuclear Reactions, Joint Institute for Nuclear

In the present review, we demonstrate the possibility of using irradiation of samples with high-energy gas ions (Kr, Xe, and Bi) for introduction and modification of Si nanocrystals in Si:SiO2 and NC:SiO2 layers, and also for formation of ordered arrays of Si nanocrystals embedded in a SiO2 matrix. Due to the small size of nanocrystals synthesized in our samples (typically, ≤ 6 - 8 nm), in describing our results we will also use an alternative term 'quantum dots'. We will analyze the changes in the optical and electronic properties of NC systems and discuss possible applications of such systems. It should be emphasized here that a special feature of used irradiation, in contrast to conventional ion implantation, is huge ion-energy losses in the bulk of Si:SiO2 layers and simultaneous creation of a relatively low defect concentration in such layers (Antonova et al., 2006a; Toulemonde et al., 2000; Meftah et al., 1997). The factor causing a low defect concentration in examined irradiated layers is the small thickness of those layers, which is normally much smaller than the projected ion range. Hence, while in practice defects are introduced in the substrate of film systems, the energy is mainly dissipated in the top Si:SiO2 layer. The latter circumstance has allowed us to separate out the positive effect of ion irradiation (the huge energy losses in ion

tracks) from the negative effect (defect formation).

**2. Sample preparation and ion irradiation conditions**

**2.1 Nanocrystals in SiO2 layers with varied concentration of excess Si**

layer composition, and also of thermal treatments given to the samples.

**2.2 High-energy ion irradiation: Specific features**

The Si:SiO2 layers of thickness 400–1000 nm used in our study were sputtered from Si and SiO2 sources onto silicon substrates. In the experimental arrangement used, the SiO2 and Si sources were spaced ~100 mm apart. The large distance between the sources provided a possibility to form layers with varied concentration of excess silicon in the oxide. After annealing of samples and formation of nanocrystals, we could obtain layers with concentration and sizes of NCs varied in the lateral plane. That gave us an interesting opportunity to examine the properties of formed Si:SiO2 layers as dependent on NC concentration and sizes. The excess Si content of the silicon-rich oxide layers ranged from 4 to 94 vol.%. The composition of deposited layers was estimated from the thickness of Si and SiO2 layers obtained by individual deposition of films from the Si and SiO2 sources onto different substrates (Shamin et al., 2010). Subsequent high-temperature annealing (at 1140°C for 40 min) used for part of samples resulted in the formation of silicon oxide layers containing Si nanocrystals (NC:SiO2 layers), whose sizes only slightly depended on layer composition and whose concentration increased with coordinate along the sample by several orders of magnitude (Sa'ar et al., 2005; Antonova et al., 2006b). Experimental results are presented below as a function of excess Si concentration. This parameter is used to characterize the composition of our Si:SiO2 layers since its value was inherently defined by the evaporation procedure, whereas other parameters, such as NC concentration and sizes, were complicated functions of

For irradiation of Si:SiO2 and NC:SiO2 layers obtained as described above, Kr, Xe, or Bi ions respectively with energies 90, 130-250, and 65 - 670 MeV were used. The ion fluences ranged from 3x1011 to 3x1013 cm–2. The irradiations of our samples were performed with the use of accelerators available in the Laboratory of Nuclear Reactions, Joint Institute for Nuclear Research, Dubna. For the ion energies used, the ion penetration range was 15–23 μm (see Table 1); i.e. it was substantially larger than the thickness of silicon-rich oxide layers. This means that almost all ions penetrated into the substrate, and the defects generated in elastic scattering channels were also predominantly introduced in the substrate, whereas prevailing energy losses in the layers were the electron stopping losses.

With reference to SRIM data, the electron stopping losses due to ionization dEe/dx in the oxide layers were expected to vary from 7.1 (Kr), 10.2 (Xe), and 22.0 keV/nm (Bi) for the SiO2 end of the sample to 9.5 (Kr), 12.3 (Xe), and 24.3 keV/nm (Bi) for its Si end. The ionization losses of ion energy lead to ionization of atoms in the material under irradiation. In a time interval of 10-13 - 10-12 s the excitation energy was transferred into anisotropic heating of the oxide layers due to thermalization process. Using the value of ionization losses, the dependence of ion-track sizes in -SiO2 on ion energy, and the thermal spike model of (Toulemonde et al., 2000; Meftah et al., 1997), we could estimate the width of ion tracks in our samples, which turned out to be about 4-6 nm. Results for the time evolution of temperature inside ion tracks calculated by TSRIKE-02 software (Toulemonde et al., 2000) are shown in Fig.1a. TSPIKE-02 simulates the electron-lattice interaction (transfer of the ionization losses of ions into the heating of the target lattice). The variable parameter of the ion-lattice interaction is the radius of the shells surrounding individual ion trajectories in which the main portion of ion energy is lost for ionization. As it follows from Fig.1a (see also Fig.15 a), the local transient temperature in ion track cores in our experiments could reach 1700 - 2500oC, the ion-track heating time being about 10-11 - 10-10 s. It means that almost all inter-atomic bonds along ion tracks could be broken to free the atoms and make them capable of displacing from their regular lattice sites. As a result, new bonds could readily form in the material, this process leading to subsequent amorphization or crystallization of part of previously existing Si nanoparticles as well as to the formation of new NCs in the oxide layer along ion tracks.

One more effect that was taken into consideration for materials subjected to the irradiation with high-energy ions (see also Section 6) is ion-beam-induced anisotropic plastic deformations (Van Dillen et al., 2004; 2005). The strain anisotropy is imposed by the direction of the ion beam: normally, the material expands in the direction normal to the ion beam and contracts parallel to it, while maintaining their volume. According to the developed model, the strain (determined as the relative change of lattice constant) versus time goes through a maximum and then reaches a constant value. The anisotropic strain rises with ion energy and fluence. It should be emphasized here that up to fluences 1012 cm-2 the ion tracks in the material were independent of each other, whereas at higher fluences they can become overlapping.

The ionization losses of the ions and the anisotropic heating of the material are higher in Si in comparison with SiO2. The anisotropic deformations found in amorphous silicon are typically 1.5 orders of magnitude lower than the anisotropic deformations in SiO2 (Van Dillen et al., 2005; Meftah et al., 1997). This means that the anisotropic strain and the anisotropic heating should demonstrate considerable changes with increasing concentration of Si nanocrystals embedded in the SiO2 layer. In the case of variable NC concentration we have to expect a strong anisotropic-strain effect for SiO2 layers with low Si-phase content and a more pronounced anisotropic heating effect for SiO2 layers with high excess Si phase content.

Si Nanocrystal Arrays Created in SiO2 Matrix by High-Energy Ion Bombardment 157

In the initial (non-irradiated) layers, the spherically shaped Si nanocrystals were distributed randomly over the insulator matrix; the NC spheres were 3–5 nm in diameter for excess Si contents of Si:SiO2 layers 40-70 vol.%, with the atomic planes in the nanocrystals having random orientation. After ion irradiation with ionization losses over 6 keV/nm, in the oxide layers with a relatively high excess Si content (above 35%) vertically ordered NC arrays were formed along ion tracks (Fig.2) (Antonova et al., 2009 a; 2009 b). The morphology of NCs has also changed: the nanocrystals became extended along the ion tracks. As shown in Fig. 2 b, this effect is accompanied by another effect: the atomic planes in the formed NCs were aligned in the direction along ion tracks, with deviation ± 20o. It should be emphasized here that no other method of forming Si NCs allows production of vertically ordered arrays

The TEM images presented in Fig.2 were obtained for samples NC-11 and NC-21 irradiated with 90 MeV Kr ions; the excess Si content of the oxide layer in those samples was 37–45%. The Kr ion had a minimal mass and a relatively low energy. We believe that in the case of using for irradiation heavier ions with a higher energy such effects will become essentially

In irradiated Si:SiO2 structures which initially contained no NCs (sample NC-22), introduction of Si nanocrystals with a similar spatial and size distribution was revealed after

Fig. 2. Cross-sectional TEM and HRTEM images of NC:SiO2 layers in samples NC-21 irradiated with 90 MeV Kr ions to a fluence of 1012 cm-2. The Si phase contents of the samples were about 37 (a) and 73 vol.% (b). The arrows show the direction of ion tracks. Note that in the irradiated sample (b) the atomic planes inside nanocrystallites are aligned

**3.2 Amorphization and crystallization processes studied by Raman spectroscopy**

Figure 3 a, b compares Raman spectra taken from initial and irradiated samples in areas with two different Si-content values, 51% and 62% (Antonova et al., 2009 a; 2010 a). The

**3. Structural modification of Si nanocrystals embedded in SiO2 matrix**

**3.1 Ordered nanocrystal arrays in ion-modified layers**

of nanocrystals with identically oriented atomic planes.

additional high-temperature thermal treatments given to the samples.

more pronounced.

along ion tracks (with deviation ± 20о).

Table 1. Parameters of structures used, and annealing and irradiation regimes. Rp is the projected ion range in Si calculated by SRIM. Eex values are the ionization losses of the ions in SiO2.

Fig. 1. Simulated evolution of temperature in time in -SiO2 (quartz) samples for increasing radial distances from the ion track axis. The simulations were performed using the TSPIKE-02 software (Toulemonde et al., 2000). For times shorter than 10-13 sec the temperature should be considered as quasi-temperature.

## **3. Structural modification of Si nanocrystals embedded in SiO2 matrix**

## **3.1 Ordered nanocrystal arrays in ion-modified layers**

156 Ion Implantation

NC-11a, 700 1140оС, 40 min Kr, 90 1х1012 16 7.1 800оС, 20 min

NC-23a, 1000 1140оС, 40 min Xe, 130 1х1012 17 10.2 800оС, 20 min NC-24, 1000 1140оС, 40 min Xe, 130 1х1013 17 10.2 800оС, 20 min

NC-31a, 400 1140оС, 40 min Kr, 250 1х1012 33.1 10.2 800оС, 20 min

NC-32a, 400 1140оС, 40 min Kr, 250 2х1013 33.1 10.2 800оС, 20 min NC-33, 400 1140оС, 40 min Bi, 65 1х1012 10 8.9 800оС, 20 min NC-34, 400 1140оС, 40 min Bi, 130 1х1012 16 13.2 800оС, 20 min NC-35, 400 1140оС, 40 min Bi, 670 3х1011 43 18.0 800оС, 20 min NC-36, 400 1140оС, 40 min Bi, 670 1х1012 43 18.0 800оС, 20 min NC-37, 400 1140оС, 40 min Bi, 670 8х1012 43 18.0 800оС, 20 min NC-38, 400 1140оС, 40 min Bi, 670 3х1013 43 18.0 800оС, 20 min Table 1. Parameters of structures used, and annealing and irradiation regimes. Rp is the projected ion range in Si calculated by SRIM. Eex values are the ionization losses of the

Fig. 1. Simulated evolution of temperature in time in -SiO2 (quartz) samples for increasing radial distances from the ion track axis. The simulations were performed using the TSPIKE-02 software (Toulemonde et al., 2000). For times shorter than 10-13 sec the temperature

should be considered as quasi-temperature.

NC-21, 1000 1140оС, 40 min Kr, 90 1х1012 16 7.1 - NC-22, 1000 - Xe, 130 1х1012 17 10.2 1140оС, 40 min

Fluence, cm-2

Rp in Si, mm

Eex, keV/nm Annealing after irradiation

Ion energy, MeV

NC-11, 700 1140оС, 40 min Kr, 90 1х1012 16 7.1

NC-23, 1000 1140оС, 40 min Xe, 130 1х1012 17 10.2

NC-31, 400 1140оС, 40 min Kr, 250 1х1012 33.1 10.2

NC-32, 400 1140оС, 40 min Kr, 250 2х1013 33.1 10.2

Sample, thickness of SiOx layer, nm

ions in SiO2.

Annealing regime before irradiation

NC-10, 700 1140оС, 40 min

NC-20, 1000 1140оС, 40 min

NC-30, 400 1140оС, 40 min

In the initial (non-irradiated) layers, the spherically shaped Si nanocrystals were distributed randomly over the insulator matrix; the NC spheres were 3–5 nm in diameter for excess Si contents of Si:SiO2 layers 40-70 vol.%, with the atomic planes in the nanocrystals having random orientation. After ion irradiation with ionization losses over 6 keV/nm, in the oxide layers with a relatively high excess Si content (above 35%) vertically ordered NC arrays were formed along ion tracks (Fig.2) (Antonova et al., 2009 a; 2009 b). The morphology of NCs has also changed: the nanocrystals became extended along the ion tracks. As shown in Fig. 2 b, this effect is accompanied by another effect: the atomic planes in the formed NCs were aligned in the direction along ion tracks, with deviation ± 20o. It should be emphasized here that no other method of forming Si NCs allows production of vertically ordered arrays of nanocrystals with identically oriented atomic planes.

The TEM images presented in Fig.2 were obtained for samples NC-11 and NC-21 irradiated with 90 MeV Kr ions; the excess Si content of the oxide layer in those samples was 37–45%. The Kr ion had a minimal mass and a relatively low energy. We believe that in the case of using for irradiation heavier ions with a higher energy such effects will become essentially more pronounced.

In irradiated Si:SiO2 structures which initially contained no NCs (sample NC-22), introduction of Si nanocrystals with a similar spatial and size distribution was revealed after additional high-temperature thermal treatments given to the samples.

Fig. 2. Cross-sectional TEM and HRTEM images of NC:SiO2 layers in samples NC-21 irradiated with 90 MeV Kr ions to a fluence of 1012 cm-2. The Si phase contents of the samples were about 37 (a) and 73 vol.% (b). The arrows show the direction of ion tracks. Note that in the irradiated sample (b) the atomic planes inside nanocrystallites are aligned along ion tracks (with deviation ± 20о).

## **3.2 Amorphization and crystallization processes studied by Raman spectroscopy**

Figure 3 a, b compares Raman spectra taken from initial and irradiated samples in areas with two different Si-content values, 51% and 62% (Antonova et al., 2009 a; 2010 a). The

Si Nanocrystal Arrays Created in SiO2 Matrix by High-Energy Ion Bombardment 159

Samples NC-36 and NC-37 to be discussed in the present section were irradiated with Bi ions of about 670 MeV energy (the projected ion range 43 nm) to fluences in the range 1012 – 1013 cm–2. Those ions, having a maximal mass and energy, were expected to provide the most significant effect on Si NCs. The specific ionization losses of 670 MeV Bi ions in the SiO2 layers were 24 keV/nm, which value far exceeded the threshold ionization energy (~2 keV/nm) required for the formation of latent tracks in SiO2 (Toulemonde et al., 2000). Prior to performing Q-DLTS measurements, the radiation defects introduced in the substrate were

As a rule, Q-DLTS spectra of non-irradiated and irradiated samples involved a single peak. In the temperature range 80 to 300 K, this peak features two slopes in the Arrhenius dependence (Fig. 4a). The activation energies determined from the measurements made at several points on a sample with different compositions of the NC:SiO2 layer are listed in Table 2. A hypothesis was put forward that the levels detected by Q-DLTS in the NC:SiO2 layers were quantum confinement levels in Si nanocrystals. In the latter case, the energy interval between the levels could be expected to be dependent on NC sizes. Assuming that ΔEi = En=2-En=1 (here, En=2 and En=1 are the energies of the first excited and ground state of an electron in the quantum dot), we were able to evaluate the typical NC size value W from the theoretical relation between En and W. Starting with the effective mass approximation (Moskalenko et al., 2007), we have obtained, from the theoretically estimated values of ΔEi, the W values given in the last column of Table 2. Trying then to approximate the dependence En(W) with the simpler well-known expression (see, for

> 2 2 <sup>2</sup> , <sup>2</sup>

where En is the energy of the n-th confined level and m is the electron effective mass in crystalline Si (0.26 in units of the free electron mass), we have obtained *W* values given in the last column of the table as bracketed values. We see that the estimated values of *W* fall into the 3-6 nm range for the reference (non-irradiated) sample NC-30. These values of *W* show a good correlation with TEM data. Q-DLTS measurements yield practically identical sets of levels for various points measured along samples, this finding being in turn indicative of an identical set of NC sizes. As it follows from Table 2, the irradiation brings about a decrease of NC sizes. A further decrease in NC sizes was observed in samples irradiated to higher fluences. In the latter case, formation of smaller Si nanocrystals might proceed due to three different reasons. The first reason is the formation of new small Si nanocrystals in the samples under irradiation. The second reason is possible decrease in the sizes of already existing NCs, which could be a result, for instance, of oxidation of the nanocrystals. It should be noted here that dissolution of metallic nanoinclusions (in particular, Au inclusions) in samples irradiated with high-energy ions was observed by Rizza et al. (Rizza et al., 2007 a; 2007 b). The third reason, operating in a situation in which the shape of NCs could be altered by ion irradiation from spherical to ellipsoidal one, extended along ion tracks, is that the energies of quantum confinement levels were primarily defined by the smallest NC size. In the latter case, a decrease in NC sizes could be indicative of an elongation of nanocrystals and of a decrease of their effective diameter. Photoluminescence (PL) measurements performed on the same structures yields for the blue

*E n <sup>n</sup> m W* (1)

removed from it giving the samples an additional anneal at 800°C during 20 min.

example, Gaponenko, 1998),

Raman line due to the single-crystal Si substrate is detected at 521 cm-1. After the samples were irradiated with Kr or Bi ions, the amorphous 480 cm-1 line has increased significantly. In the case of Bi irradiation, an NC-related line at 500-520 cm-1 was also observed in the spectra. To demonstrate this effect, we have plotted in Fig. 2a the difference between the spectrum of a sample irradiated with Bi ions and the spectrum of the same sample before irradiation. An NC-related peak at 500-520 cm-1 is also observed in the differential spectrum; this fact points to the formation of new NCs in samples subjected to irradiation. An increase in the irradiation dose (> 1012 cm-2) and a decrease in ion energy and/or ion mass lead to domination of the amorphization process. Raman measurements performed on samples with higher Si contents (> 70 vol.%) have shown that, in the latter case, only amorphization of Si NCs was observed. Following an additional anneal at 800oC for 20 min, no substantial changes were observed in the Raman spectra.

It was found that irradiation of Si:SiO2 structures that initially contained no previously formed NCs also brought about the formation of new Si nanocrystals in the course of irradiation. This effect was observed practically for all ions used in the present study. In this case, the formation of new nanocrystals also prevails at relatively low Si contents of Si:SiO2 layers (below 50–60%), whereas the formation of amorphous inclusions prevails at higher Si contents. An increase in irradiation dose (over 1012 cm–2) and a decrease in the specific ionization losses of the ions yield a net increase in the amount of optically active nanocrystals available in the samples (Antonova et al., 2010 b).

#### **3.3 Evolution of NC sizes under ion irradiation**

It was also shown that a study of Si NCs by means of charge deep-level transient spectroscopy (Q-DLTS) makes it possible to observe the capture of charge carriers by, and their emission from, dimensional quantization levels in NCs located close to the SiO2/Si interface (Antonova et al., 2009 c; 2011). The Q-DLTS method was also employed to investigate into variations of energy barriers involved in the emission process of charge carriers from NC levels in NC:SiO2 layers modified with high-energy ion irradiation.

Fig. 3. Raman spectra of NC-SiO2 layers on Si substrates for initial (non-irradiated) samples with (a) 51% and (b) 62% of the Si phase and for samples irradiated with Kr (NC-32) and Bi ions (NC-36). In order to demonstrate the added intensity around the "amorphous" 480 cm-1 line and the NC-related 500-520 cm-1 line due to Bi+ irradiation, in the insert to the figure we have plotted the difference between the spectra measured after and before the irradiation of the sample with 670 MeV Bi+ ions (NC-30 and NC-36).

Raman line due to the single-crystal Si substrate is detected at 521 cm-1. After the samples were irradiated with Kr or Bi ions, the amorphous 480 cm-1 line has increased significantly. In the case of Bi irradiation, an NC-related line at 500-520 cm-1 was also observed in the spectra. To demonstrate this effect, we have plotted in Fig. 2a the difference between the spectrum of a sample irradiated with Bi ions and the spectrum of the same sample before irradiation. An NC-related peak at 500-520 cm-1 is also observed in the differential spectrum; this fact points to the formation of new NCs in samples subjected to irradiation. An increase in the irradiation dose (> 1012 cm-2) and a decrease in ion energy and/or ion mass lead to domination of the amorphization process. Raman measurements performed on samples with higher Si contents (> 70 vol.%) have shown that, in the latter case, only amorphization of Si NCs was observed. Following an additional anneal at 800oC for 20 min, no substantial

It was found that irradiation of Si:SiO2 structures that initially contained no previously formed NCs also brought about the formation of new Si nanocrystals in the course of irradiation. This effect was observed practically for all ions used in the present study. In this case, the formation of new nanocrystals also prevails at relatively low Si contents of Si:SiO2 layers (below 50–60%), whereas the formation of amorphous inclusions prevails at higher Si contents. An increase in irradiation dose (over 1012 cm–2) and a decrease in the specific ionization losses of the ions yield a net increase in the amount of optically active

It was also shown that a study of Si NCs by means of charge deep-level transient spectroscopy (Q-DLTS) makes it possible to observe the capture of charge carriers by, and their emission from, dimensional quantization levels in NCs located close to the SiO2/Si interface (Antonova et al., 2009 c; 2011). The Q-DLTS method was also employed to investigate into variations of energy barriers involved in the emission process of charge

carriers from NC levels in NC:SiO2 layers modified with high-energy ion irradiation.