SiN dielectric stacks by low energy ion beam synthesis

Thin Solid Films 543 (2013) 94–99

Contents lists available at ScienceDirect

Thin Solid Films journal homepage: www.elsevier.com/locate/tsf

Ge nanocrystals in HfO2/SiN dielectric stacks by low energy ion beam synthesis M. Carrada a, b,⁎, B.S. Sahu b, 1, C. Bonafos a, F. Gloux a, J. Groenen a, D. Muller b, A. Slaoui b a b

CEMES/CNRS, 29 rue J. Marvig, 31055 Toulouse, France InESS, CNRS, 23 rue du Loess, 67037 Strasbourg, France

a r t i c l e

i n f o

Available online 5 March 2013 Keywords: Ge-NCs Ge nanocrystals Ion beam synthesis HfO2/Si3N4 stacks Non-volatile-memories

a b s t r a c t Germanium nanocrystals (Ge-NCs) have been obtained by low energy ion beam synthesis in a SiNx/HfO2 stack layer. The effect of the Ge implanted dose variations on structural characteristics (size, position, chemical bonding) of Ge-NCs have been investigated by Transmission Electron Microscopy and Raman spectroscopy. Our results show that several processes (damage, diffusion, oxidation …) that depend on the Ge implanted dose, take place during the synthesis and complicate the expected behavior of the ion beam synthesized system. However, significant memory windows with good retention properties have been observed in these stack structures, indicating their feasibility for low operating voltage, non-volatile memory devices. © 2013 Elsevier B.V. All rights reserved.

1. Introduction Semiconductor nanocrystals (NC), in particular silicon and germanium, have been extensively investigated in the last decade for their application to non-volatile memories (NVM) [1,2]. Following the different research groups, Si [3,4] or Ge [5,6] nanocrystals have been preferred, each of them presenting advantages and disadvantages. For example when embedded in silicon dioxide, controlling Si-NCs is easier than Ge-NCs because they are less exposed to oxidation and diffusion and hence, their implementation for industrial application is generally preferred. But using Ge-NCs instead of Si-NCs has been expected to increase the retention time in memory devices as the smaller band gap of Ge should provide a deeper potential well for the stored charges. In reality, only the hole retention time is higher for Ge-NC than for Si-NCs, as shown by predictive simulations [7] and experiments [8–10], while the barrier height for electrons is almost the same for Ge and Si. As the retention time for holes in Ge is much higher than that in Si, direct tunneling is in principle allowed [7,11] and hence, writing/erasing cycles should be possible without the tunnel oxide degradation, which is a strong argument for preferring Ge-NCs over Si-NCs for memory applications. Moreover, the advantage of using Ge-NC based memories rather than Si-NC ones is supported by the experimental results obtained by King et al. [12] in terms of writing/erasing times and by Kanoun et al. in terms of data retention [13]. However, other authors [14] have demonstrated the presence of high level concentration of charge trap states at the Ge-NC/SiO2 matrix interface, which would enhance trap assisted out-tunneling and charge losses, concluding that Ge-NCs would not ⁎ Corresponding author at: CEMES/CNRS, 29 rue J. Marvig, 31055 Toulouse, France. Tel.: + 33 562257963. E-mail address: [email protected] (M. Carrada). 1 Now at Department of Physics, KIIT University, 751024 BBSR, Odisha, India. 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2013.02.113

provide advantages on Si-NCs for memory applications. On the contrary, other groups [15] consider the presence of these traps advantageous and propose a trap level engineering through the tunnel oxide choice and Ge-NCs size tuning for long retention memory devices. To conclude this short overview on the debate Ge vs. Si, it is clear that the advantage of using Ge-NCs with respect to Si-NCs is still a controverted question and for the moment no direct comparison of equivalent memory structures containing Ge-NCs and Si-NCs has been made in the literature. Finally, it must be noticed that all the exposed considerations are referred to Ge or Si-NCs embedded in SiO2 gate dielectrics. For non-volatile memories, the replacement of the silicon oxide by a gate dielectric with higher permittivity (high-k) has been recently proposed in order to obtain very thin equivalent oxide thickness (EOT) and to improve data retention and programming speed in nanocrystal (Ge or Si) based non-volatile memories. Indeed, the use of high-k dielectric allows for higher physical thicknesses of the tunnel oxide, which correspond to thinner EOT. In this context, both Ge and Si-NCs have a high potential for their application to NVM, which structure can be modulated to optimize their properties. Ge-NC based NVM using high permittivity dielectrics including HfO2 [16], Al2O3 [17], ZrO2 [18], Lu2O3 [19] or stack of them (ZrO2/ Al2O3…) [20] have been reported. HfO2 seems to be the most promising gate dielectric for NVM application because of its very high dielectric constant value and proper band offset, and tunnel barrier asymmetry, allowing to improve the trade-off between programming speed and non-volatility. Indeed, in the case of a symmetric tunnel barrier the probability for tunneling in and out of a NC is the same, and hence, long retention times are connected with long writing times. On the contrary, band asymmetry with respect to the Si substrate is useful to get long retention times and very short programming times [21,22]. Consequently, the trade-off between programming speed and non-volatility can be further improved by

M. Carrada et al. / Thin Solid Films 543 (2013) 94–99

increasing the asymmetry of the tunnel barrier [23,24], for example by using stack tunnel layers. Following this route, in this work we have studied an original memory structure in which Ge-NCs have been obtained by low energy ion beam synthesis (LE-IBS) in thin HfO2/Si3N4 stack dielectric. A similar structure with Ge-NCs and stacked HfO2/SiO2 tunnel dielectrics obtained by electron beam evaporation has been studied by Wang and Lu [25]. However, they have employed relatively thicker dielectric films with respect to ours. Our system takes advantage of the HfO2 ideal physical properties as tunnel oxide, while Ge-NCs are used as charge storage nodes. Moreover, we have used silicon nitride as a control dielectric, because of his double advantage: indeed, it increases the barrier asymmetry and it is also an excellent diffusion barrier preventing the Ge atoms to go towards Si substrate and/or sample free surface [26]. The method we have used, LE-IBS is very interesting for synthesizing nanocrystals embedded in dielectric matrix. In fact, two dimensional arrays of NCs with an extremely narrow size distribution and spatially organized in a plane parallel to the dielectric/Si substrate interface can be obtained. The NC layer position with respect to the electrodes and the NC population characteristics (density and size) can be manipulated by fabrication parameters (implantation energy, implanted dose, annealing conditions…) with sub-nanometric precision [27]. 1.1. Experimental details The systems under investigation are obtained by the following way: before ion implantation, 1.2 nm of SiO2 was thermally grown on p-type Si (100) substrates (resistivity 1–10 Ω·cm). Subsequently, 4.7 nm of HfO2 was deposited by metal organic chemical vapor deposition technique. The top SiN layer with a thickness of about 12 nm was then deposited with Electron Cyclotron Resonance — Plasma Enhanced Chemical Vapor Deposition (ECR-PECVD) method under a flow of SiH4 and N2 (instead of NH3) to minimize the H content in the films. Ion implantation in these stack layers were carried out with 74Ge + ions using GeH4 gas source for the extraction of Ge. The Ge + ion implantation was carried out at a constant energy value of 5 keV, which was found to be the optimum condition in previous studies [28]. The implanted doses were the following: 1 × 10 16 cm−2, 1.5 × 1016 cm−2 and 2 × 10 16 cm−2, corresponding to the samples called D1, D1.5 and D2, respectively. The post-implanted samples were subjected to conventional furnace annealing at 800 °C in highly pure dry N2 for 30 min in order to obtain the Ge-NC formation. The formation and evolution of Ge-NCs as a function of the Ge implanted dose have been investigated using high-resolution electron microscopy (HREM) and Energy Filtered Transmission Electron Microscopy (EFTEM) on cross-sectional specimens. The EFTEM images are formed with the electrons that are selected by a slit placed in the energy-dispersive plane of the spectrometer with a width of 4 eV centered at an energy position of 16 eV (Ge-plasmon energy). The samples for TEM observations were prepared by mechanical polishing and ion milling using the standard procedure. HREM and EFTEM images were taken using a field emission TEM (FEI Tecnai™ F20 operating at 200 kV) equipped with a spherical aberration corrector. Raman spectroscopy was also used in order to optically evidence the presence of Ge-NCs. Raman scattering measurements were performed with a DILOR-XY spectrometer at room temperature, using the 514.4 nm line of an argon laser for excitation. The atomic composition of the layers before and after the thermal annealing has been analyzed by Rutherford Backscattering Spectroscopy (RBS) using He+ particles with 2 MeV energy obtained with a Van De Graaff accelerator. The RBS spectra were obtained in classical configuration: the sample inclination with respect to the beam was 10° and the diffusion angle was 160°. Finally, metal–insulator–semiconductor (MIS) memory capacitor structures were fabricated from the samples by evaporating Al electrodes with

95

0.8-mm diameter with a shadow mask and Al rear side contact after scratching the back surface. Capacitance–voltage (C–V) and conductance–voltage (G–V) measurements were carried out using HP4192A impedance analyzer through a LABVIEW interface. 2. Results and discussion Fig. 1a, b, and c shows the cross-sectional HREM images of samples D1, D1.5 and D2, respectively. First of all, we can notice that for all the 3 different doses Ge nanocrystals are formed. The measured lattice fringes separation is 0.327 nm, which matches well with the Ge (111) inter planar distance in the diamond structure. When comparing the images corresponding to the 3 different samples, we observe that when increasing the dose the Ge-NCs are positioned at a larger distance from the HfO2 layer. Indeed, the injection SiN (between HfO2 and Ge-NCs) layer thicknesses are respectively 3.1 nm, 5.7 nm and 8.6 nm for samples D1, D1.5 and D2. At the same time a swelling of the silicon nitride matrix is observed, whose values increase when the implanted dose increases. Indeed the total thickness of the silicon nitride layer increases from 11.9 nm for sample D1 to 15.7 nm for sample D1.5 and to 18.2 nm for sample D2. This swelling is not in agreement with the simulation performed by the TRIDYN code [29], which predicts a contraction of the SiN layer after Ge ion implantation instead of swelling as a result of Ge atoms addition and surface sputtering. On the other hand, there is no significant increase of the HfO2 layer thickness while the interface layer (IL) thickness is 2 nm for the lowest doses and 0.8 for the higher dose. The control dielectric thicknesses (between the Ge-NCs and the free surface of the sample) are respectively 3 nm, 5.5 nm and 1.4 nm for samples D1, D1.5 and D2. When observing the Ge-NCs in EF-TEM images for samples D1 and D2 (Fig. 2a and b) we can notice than in the first case only one band of NCs is formed, with high density, and in the second case two layers of Ge-NCs are observed, which appears less dense. In principle, an EFTEM plan view study would be helpful to compare Ge-NC density values for the three doses and to confirm that the shift of Ge towards the surface is accompanied by a loss of Ge as indicated by Raman spectroscopy and RBS. However, the PV study of such stack layers has been difficult and the results are not useful for quantitative comparison due to the layers' small thicknesses. Indeed, HfO2 and Ge plasmon signature are superposed and if we want to be sure to image by EFTEM only the Ge-NCs and not also the HfO2 grains, we should obtain thin areas avoiding superimposition of SiN layer with Ge NCs and the HfO2 grains below. This has been impossible in our case because in these structures the Ge-NCs are too close to the HfO2 layer, at the scale of sample preparation for TEM. All the obtained results reveal an unexpected evolution of the system when the Ge implanted dose increases. Indeed, due to the low Ge diffusion in silicon nitride, the Ge-NCs should be formed at the ion implantation projected range (Rp = 8.3 nm, maximum of the implanted atoms concentration in the layer), and their position should be only determined by the implantation energy, which is constant in this case. But as we have described, when the implanted dose increases, a shift of the Ge-NCs to the surface is observed, accompanied by a decrease of Ge-NC density, and a strong and increasing swelling of the silicon nitride layer. As already mentioned, the measured swelling values cannot be explained by Ge atoms addition due to ion implantation, even if we neglect the sputtering effects. In order to better understand this behavior, we have performed Raman Spectroscopy and Rutherford Backscattering Spectroscopy (RBS) analysis. Raman spectra are reported in Fig. 3. In addition to the huge silicon substrate signal, weak features are observed at about 300 cm −1. This spectral range is characteristic for Ge optical phonons. The sharp peak observed in the upper spectrum clearly indicates that Ge is crystallized: the sample with 10 16 cm −2 dose and annealed at

96

M. Carrada et al. / Thin Solid Films 543 (2013) 94–99

Fig. 1. HREM cross sectional images of samples (a) D1 (1 × 1016 Ge cm−2), (b) D1.5 (1.5 × 1016 Ge cm−2) and (c) D2 (2 × 1016 Ge cm−2). When the Ge implanted dose increases, the total SiN thickness increases (swelling), and a shift of the Ge-NCs toward the sample free surface is observed. Only one Ge-NC plane is observed for the lower doses (a and b), while for the highest dose the Ge-NCs are distributed in two different planes (c).

950 °C does contain Ge-NCs. For the sample with the same dose but lower annealing temperature (800 °C) only a weak feature is observed. For the highest dose (1.5 × 10 16 cm −2), no clear features are present whatever the annealing temperature. The absence of sharp Ge-related peaks could be consistent with a weak amount of crystallized Ge or significant oxidation. However, as the HREM study revealed crystallographic distances of Ge (3.27 Å) and not of GeO2 (3.43 Å) the hypothesis of a loss of Ge\Ge bonds when increasing the Ge implantation dose is more consistent. The Ge atoms distribution obtained by RBS before and after the thermal annealing is shown in Fig. 4. These profiles shows that the implanted Ge atoms stay in the silicon nitride layer, without significant diffusion at Si/HfO2 interface, as required for NVM good performances. The layers composition as well as the Ge content obtained by RBS analysis is summarized in the Table 1. As shown in Table 1, the Ge dose remaining in the layers after the 800 °C, 30 min thermal annealing is 0.94 × 10 16 cm − 2; 0.7 × 1016 cm−2 and 1.3 × 10 16 cm−2, respectively for samples D1, D1.5 and D2 whose nominal doses are respectively 1 × 1016 cm−2, 1.5 × 1016 cm−2 and 2 × 1016 cm−2. This means that only for the lowest implanted dose there is almost no loss of Ge atoms, while for the two highest doses more than 40% of the implanted Ge atoms are lost. However, the amount of Ge in the matrix is sufficient to obtain Ge-NCs. Another important result obtained by the RBS analysis is

the presence of a considerable amount of oxygen in the silicon nitride matrix. This oxygen is already present after the Ge implantation, before the thermal annealing, and seems to be uniformly distributed from the free surface to the Ge implanted silicon nitride. The composition of the matrix is found to be close to Si1O1N1 for the lowest implanted doses and close to Si1O2N1 for the highest one, instead of Si3N4 before the ion implantation. On the contrary, the silicon nitride zone close to the HfO2 layer, and without Ge atoms, contains only traces of oxygen atoms. In addition, we have to mention that previous Elastic Recoil Detection Analysis in similar samples have shown that our ECR-PECVD deposited silicon nitride layers contains between 15 at.% and 30 at.% of hydrogen. The comparison of all the obtained results leads to a clearer interpretation of the abnormal behavior of our system as a function of the implanted dose. The shift of Ge-NCs toward the surface and the Ge loss, as well as the abnormal swelling can be explained by taking into account the SiN matrix oxidation. Indeed, the RBS measurement have shown that in the Ge implanted zone the silicon nitride matrix has turned to silicon-oxy-nitride, due to a large amount of oxygen penetration. This oxygen penetration occurs after the implantation and before the annealing. In fact, when the sample is taken out from the under vacuum implanter chamber, the water from humidity air can easily penetrate the layers which have been strongly damaged during the ion implantation. In the past, it has been shown for Si, Ge and Sn implanted in SiO2 that oxygen easily penetrates the layers

Fig. 2. Energy filtered (EFTEM) cross sectional images of samples (a) D1 (1 × 1016 Ge cm−2), with one Ge-NC band appearing in white and (b) D2 (2 × 1016 Ge cm−2) with two Ge-NC bands located very close to the sample surface.

M. Carrada et al. / Thin Solid Films 543 (2013) 94–99

16

400

950

200

Table 1 Sample composition as obtained by RBS.

-2

16

1.5 10 (cm )

800

-2

10 (cm )

Sample

Matrix composition

Total Ge content

o

D1 as implanted D1 annealed D1.5 as implanted D1.5 annealed D2 as implanted D2 annealed

Si1O1N1 Si1O1N1 Si1O1N1 Si1O1N1 Si1O2N1 Si1O2N1

1 0.94 1.2 0.72 1.34 1.3

Si substrate

Raman Intensity (arb. units)

600

97

Ge NCs

o

300

400

500

600

Raman shift (cm-1) Fig. 3. Raman spectra of samples D1 (red line, Ge implanted dose 1 × 1016 cm−2) and D1.5 (black line, Ge implanted dose 1.5 × 1016 cm−2) annealed at 800 °C and 950 °C under N2 during 30 min in a conventional furnace. Ge\Ge bonds are detectable only for the lowest implanted dose.

which are damaged by the ion implantation, and that its amount can be as high as 10 at.% [30]. Moreover, the oxygen amount in the layers increases with the implanted dose, i.e. with damage. The same authors have put in evidence that the high amounts of incorporated hydrogen and oxygen can interfere significantly with the precipitation and Ostwald ripening of nanocrystals during IBS, resulting in unexpected nanocrystal distributions, especially in very shallow implanted layers. Indeed, chemical reaction of the implanted species (Ge in our case) with hydrogen and oxygen from H2O (or OH − and H +, H3+ molecules) takes place during the annealing. This increases the solubility and diffusivity of Ge by forming GeHx complexes, leading to Ge loss by outward diffusion of volatile Ge compounds as GeH4 or GeO [31,32]. Moreover, it has been shown by Eugène et al. [33] that during the annealing, the presence of a small amount of Ge strongly enhances the rate of Si oxidation, and that for Ge contents b 50%, Si oxidation rate increases with Ge content. All these considerations are in agreement with the behavior of our system, in which the oxygen content increases when the Ge implanted dose increases. The enhanced Si oxidation rate when the Ge content increases is consistent with the increasing swelling due to the higher oxidation of the Si atoms in the Si3N4 matrix. Indeed, the oxidation of Si leads to volume enhancement, as shown by Deal and Groove's classical oxidation theory [34], for which the oxidation of 0.44 nm of Si leads to the formation of 1 nm of SiO2 due to the molecular densities ratio. In our case the oxidation of Si occurs within the silicon nitride matrix. However, if we consider the observed thickness (volume)

× × × × × ×

1016 1016 1016 1016 1016 1016

cm−2 cm−2 cm−2 cm−2 cm−2 cm−2

increasing, we can argue that Si is oxidized while N content is constant. Indeed, if we compare SiO2 molecular density (2.2 × 1022 at cm−3) and Si3N4 density (1.3 × 1022 at cm−3) we easily conclude that a transformation of Si3N4 in SiON by substitution of nitrogen with oxygen atoms cannot explain the measured swelling values. In fact, silicon oxide density is higher than silicon nitride, and in the case of a simple N by O substitution we expect a contraction of the layer, even assuming an intermediate density value for silicon oxy-nitride. An increasing oxygen content in the layer when the Ge implanted dose increases, is not only consistent with the higher values of the layers thickness, but also explains on one hand the Ge-NCs moving towards the surface, due to the enhanced Ge diffusion which increases with the oxygen content, and on the other hand, the Ge atom loss, due to the higher formation rate for volatile GeO. The electrical properties of these structures have also been investigated and can be understood following their structural characteristics. Fig. 5 shows the C–V curves obtained at different frequencies for samples (a) D1, (b) D1.5, and (c) D2, respectively. For the lower implanted doses (sample D1 and D1.5), larger memory windows were observed. In particular, for sample (D1.5) corresponding to the intermediate implanted Ge dose, flat band shift values of 3.95 V and 4.53 V have been obtained for sweep voltage of ±6 V and ±7 V, respectively. Moreover, the C–V curves for samples D1 and D1.5 are almost frequency independent, indicating that the charge is effectively stored in the Ge-NCs and interface defects as well as and bulk defects can be negligible. On the contrary, sample D2, implanted with the highest dose has a reduced memory window (ΔVfb ~ 1 V), as expected from its structural characteristics. Indeed, in this case Ge-NCs are located very close to the free surface of the sample, the control dielectric is very thin, and there is no significant insulation from the gate electrode. Moreover, lateral charge losses are expected due to the superimposing of two Ge-NCs bands. The charge retention has been studied (Fig. 6) for samples D1 and D1.5 which have large memory windows. In both cases we have obtained good retention properties, and for sample D1.5, implanted with 1.5 × 10 16Ge·cm −2, we have observed about 36% of charge loss after a waiting time of 10 4 s. An extrapolated value of 1.06 V is found for the memory window after 10 years, as requested for the commercial memory devices, which is very interesting for the application of these structures to NVM.

Fig. 4. RBS analysis of samples (a) D1 (1 × 1016 Ge cm−2), (b) D1.5 (1.5 × 1016 Ge cm−2) and (c) D2 (2 × 1016 Ge cm−2), before and after annealing at 800 °C for 30 min under N2.

98

M. Carrada et al. / Thin Solid Films 543 (2013) 94–99

Fig. 5. C(V) curves obtained at different frequencies, 10 kHz (blue), 100 kHz and 500 kHz (black), for the samples D1 (1 × 1016 Ge cm−2) (a), D1.5 (1.5 × 1016 Ge cm−2) (b) and D2 (2 × 1016 Ge cm−2) (c). The memory windows (variation of flat band energy ΔVfb) are very large for samples D1 (ΔVfb ~ 2.92 V at ±6 V) and D1.5 (ΔVfb ~ 3.95 V at ±6 V and ΔVfb ~ 4.53 V (±7 V), while for sample D2 the memory effect is very low (ΔVfb ~ 0.86 V at ±6 V) and ΔVfb ~ 1.01 V at ±7 V).

Sample D1.5, implanted with the intermediate dose shows the better results in terms of both charge storage and retention times. This is probably due to a better compromise in terms of distances to the electrodes and defect passivation. Indeed, its thicker tunnel and control oxides allow less leakage and partially justify the better electrical characteristics, in particular with respect to sample D2 for which the Ge-NCs are too close to the free surface of the sample. Considering sample D1, even if the tunnel and control oxides are thinner, their thickness is in principle sufficient to prevent charge leakage. For this reason, we suppose that a better passivation is obtained for sample D1.5 with respect to sample D1, due to the highest oxygen content, which on one hand passivates the Si related defect in the silicon nitride matrix [35] and on the other hand contribute to the elimination of the single Ge atoms not contained in Ge-NCs via the formation of the volatile GeO.

3. Conclusions In this work we have synthesized germanium nanocrystals by low energy ion beam synthesis (LE-IBS) in Si3N4/HfO2 stack layers. We have studied the effect of the implanted Ge dose on the structural properties of this system by using Rutherford Backscattering Spectroscopy, Transmission Electron Microscopy (HREM, EFTEM) and Raman Spectroscopy. We have observed an unexpected evolution of the system due to several processes (damage, diffusion, oxidation …) which takes place during the synthesis depending on the Ge implanted dose. In particular, the Ge-NCs position wasn't determined, as expected, by the implantation energy but it was found to be dependent on the implantation dose. Indeed, when the implanted dose increases, a shift of the Ge-NCs to the surface was observed, accompanied by a decrease of Ge-NC density. At the same time, a strong swelling of the silicon nitride layer was observed, which abnormal value couldn't be explained by Ge atoms addition. Both the shift of Ge-NCs toward the surface and the Ge loss, as well as the abnormal swelling have been explained by taking into account the SiN matrix oxidation, which takes place after the implantation and before the annealing due to the penetration of water from humidity air in implantation damaged layers. The penetration of oxygen and hydrogen in the layer is dependent on the damage, i.e. on the implanted dose, and additionally the oxidation rate of the Si in the Si3N4 matrix also increases with the Ge content. The transformation of Si3N4 in SiON is responsible of the observed swelling and of the enhanced Ge diffusivity resulting in Ge-NC shift toward the surface. Moreover, the humidity penetration is responsible of the Ge loss through the formation of Ge volatile compounds such as GeO and GeH4. Finally, significant memory windows with good retention properties have been observed in these stack structures, especially for the intermediate implanted dose (Ge, 1.5 × 10 16 at cm −2), indicating their feasibility for low operating voltage, non-volatile memory devices. Acknowledgments This work has been supported by the ANR Project ANR/ PNANO07-0053-NOMAD. The authors want to thank S. Lhostis from ST-Microelectronics in Crolles for the provision of the deposited high-k layers. References

Fig. 6. Flat-band voltage variation as a function of the waiting time (i.e. charge retention) for samples D1 (a) and D1.5 (b) having large memory windows. For sample D1.5 the charge loss after a waiting time of 104 s was 36%. An extrapolated value of 1.06 V is found for the memory window after 10 years.

[1] S. Tiwari, F. Rana, H. Hanafi, A. Hartstein, E.F. Crabbe, K. Chan, Appl. Phys. Lett. 68 (1996) 10. [2] H.I. Hanafi, S. Tiwari, I. Khan, IEEE Trans. Electron Devices 43 (1996) 1553. [3] C. Bonafos, H. Coffin, S. Schamm, N. Cherkashin, G. Ben Assayag, P. Dimitrakis, P. Normand, M. Carrada, V. Paillard, A. Claverie, Solid State Electron. 49 (2005) 1734.

M. Carrada et al. / Thin Solid Films 543 (2013) 94–99 [4] S. Jacob, G. Festes, S. Bodnar, R. Coppard, J. Thiery, T. Pate-Cazal, T. Pedron, B. De Salvo, L. Perniola, E. Jalaguier, F. Boulanger, S. Deleonibus, Proc. IEEE European Solid-State Device Research Conference, 2007, p. 410. [5] S. Duguay, J.J. Grob, A. Slaoui, Y. Le Gall, M. Amann-Liess, J. Appl. Phys. 97 (2007) 104330. [6] I. Berbezier, A. Karmous, A. Ronda, A. Sgarlata, A. Balzarotti, P. Castrucci, M. Scarselli, M. De Crescenzi, Appl. Phys. Lett. 89 (2006) 063122. [7] J.S. de Sousa, V.N. Freire, J.P. Leburton, Appl. Phys. Lett. 90 (2007) 223504. [8] M. Kanoun, A. Soufi, T. Baron, F. Mazen, Appl. Phys. Lett. 84 (2004) 5079. [9] V.V. Afanas'ev, A. Stesmans, L. Souriau, R. Loo, M. Meuris, Appl. Phys. Lett. 94 (2009) 172106. [10] Konchenko, Y. Nakayama, I. Matsuda, S. Hasegawa, Y. Nakamura, M. Ichikawa, Phys. Rev. B 73 (2006) 113311. [11] R. Peibst, M. Erenburg, E. Bugiel, K.R. Hofmann, J. Appl. Phys. 108 (2010) 054316. [12] Y.-C. King, T.J. King, C. Hu, IEEE Trans. Electron Devices 48 (2001) 696. [13] M. Kanoun, C. Busseret, A. Poncet, A. Souifi, T. Baron, E. Gautier, Solid State Electron. 50 (2006) 1310. [14] R. Peibst, J.S. de Sousa, K.R. Hofmann, Phys. Rev. B 82 (2010) 195415. [15] B.H. Koh, E.W.H. Kan, W.K. Chim, W.K. Choi, D.A. Antoniadis, E.A. Fitzgerald, J. Appl. Phys. 97 (2005) 124305. [16] S. Das, K. Das, R.K. Singha, A. Dhar, S.K. Ray, Appl. Phys. Lett. 91 (2007) 233118. [17] X.B. Lu, P.F. Lee, J.Y. Dai, Appl. Phys. Lett. 86 (2005) 203111. [18] S. Das, R.K. Singha, A. Dhar, S.K. Ray, A. Anopchenko, N. Daldosso, L. Pavesi, J. Appl. Phys. 110 (2011) 024310. [19] M.Y. Chan, P.S. Lee, V. Ho, H.L. Seng, J. Appl. Phys. 102 (2007) 094307. [20] Q. Wan, N.L. Zhang, W.L. Liu, C.L. Lin, T.H. Wang, Appl. Phys. Lett. 83 (2003) 138.

99

[21] D.W. Kim, T. Kim, S.K. Banerjee, IEEE Trans. Electron Devices 50 (2003) 1823. [22] J.J. Lee, X. Wang, W. Bai, N. Lu, D.L. Kwong, IEEE Trans. Electron Devices 50 (2003) 2067. [23] Y. Matsumoto, T. Hanajiri, T. Toyabe, T. Sugano, Jpn. J. Appl. Phys., Part 1 35 (1996) 1126. [24] K.K. Likharev, Appl. Phys. Lett. 73 (1998) 2137. [25] S. Wang, W. Lu, Appl. Phys. Lett. 86 (2005) 113105. [26] I.E. Tyschenko, V.A. Volodin, L. Rebohle, M. Voelskov, V. Skorupa, Semiconductors 33 (1999) 523. [27] C. Bonafos, M. Carrada, N. Cherkashin, H. Coffin, D. Chassaing, G. Ben Assayag, A. Claverie, T. Müller, K.H. Heinig, M. Perego, M. Fanciulli, P. Dimitrakis, P. Normand, Appl. Phys. 95 (2004) 5696. [28] B.S. Sahu, F. Gloux, A. Slaoui, M. Carrada, D. Muller, J. Groenen, C. Bonafos, S. Lhostis, Nanoscale Res. Lett. 6 (2011) 177. [29] W. Moller, W. Eckstein, Nucl. Inst. Methods Phys. Res. B 2 (1984) 814. [30] B. Schmidt, D. Grambole, F. Herrmann, Nucl. Inst. Methods Phys. Res. B 191 (2002) 482. [31] A. Markwitz, B. Schmidt, W. Matz, R. Grötzschel, A. Mücklich, Nucl. Instr. Meth. B 142 (1998) 338. [32] V. Beyer, J. von Borany, Phys. Rev. B 77 (2008) 014107. [33] J. Eugène, F.K. LeGoues, V.P. Kesan, S.S. Iyer, F.M. d'Heurle, Appl. Phys. Lett. 59 (1991) 78. [34] B.E. Deal, A.S. Grove, J. Appl. Phys. 36 (1965) 3770. [35] M. Carrada, A. Wellner, V. Paillard, C. Bonafos, H. Coffin, A. Claverie, Appl. Phys. Lett. 87 (2005) 251911.