J O U R N A L O F M AT E R I A L S S C I E N C E : M AT E R I A L S I N E L E C T RO N I C S 1 2 ( 2 0 0 1 ) 3 9 5 ± 4 0 1
Annealing effects on the structural and optical characteristics of electron beam deposited Ge-SeBi thin ®lms O. EL-SHAZLY 1 , M. M. HAFIZ 2 Department of Physics, Faculty of Science, UAE University, P.O. Box 17551, Al-Ain, United Arab Emirates E-mail:
[email protected] Measurements of absorbance and transmittance in Ge20 Se80 x Bix (0 x 10 at %) chalcogenide thin ®lms in the visible range at room temperature were carried out. The dependence of the optical absorption on the photon energy is described by the relation ahn B
hn E0 2 . It was found that the optical energy gap E0 decreases gradually from 1.93 to 1.205 eV with increasing Bi content up to 10 at %. The rate of the change decreases with increasing Bi content. The composition dependence of the optical energy gap is discussed on the basis of the concentration of covalent bonds formed in the chalcogenide ®lm. The decrease of optical energy gap with increasing Bi content is related to the increase of Bi±Se bonds and the decrease of Se±Se bonds. The effect of thermal annealing for different periods of time on the behavior of optical absorption of the as-deposited ®lms was investigated. The optical gap increases with increasing annealing time. The rate of change decreases with increasing annealing time, then E0 reaches a steady state. The increase in the values of the optical gap of the amorphous ®lms with heat treatment is interpreted in terms of the density of states model. The structure of the as-prepared and thermally annealed ®lms were investigated using transmission electron microscopy, energy dispersive analysis and X-ray diffraction. It was con®rmed that the as-prepared ®lms were in the amorphous state. Phase separation was observed after thermal annealing. The separated crystalline phases were identi®ed. # 2001 Kluwer Academic Publishers
1. Introduction
Chalcogenide glasses have been investigated intensively because of their promising technological applications. Chalcogenide glasses of Ge±Se alloys are interesting materials for infrared optics. They have a wide range of transparency extending from 0.6 to 30 mm. Also, these glasses have good mechanical properties such as hardness, adhesion, low internal stress and water resistance. It is possible to prepare glassy alloys in the Ge-Se system up to 42 at % germanium [1±3]. In the last decade, particular attention has been devoted to studying the electrical properties of the system Ge±Se±Bi (up to 13 at % Bi). This was due to the changing of the conduction type from p to n, which is seen to occur in the vicinity of 7 at % of Bi [4±6]. It has been pointed out that molecular phase separation might play an important role in this transition [7]. The Ge±Se±Bi system has attracted the attention of various investigators. The density of states in the 1 2
conduction band as well as that in the valence band in amorphous Ge20 Se80 x Bix thin ®lms (0 x 15 at %) has been investigated by inverse-photoemission and ultraviolet photoemission spectroscopies [8]. The composition dependence of thermal transport properties on the Ge20 Se80 x Bix system was studied by Thomas and Philip [9]. The heat treatment of chalcogenide ®lms results in a decrease of the degree of disorder and defects present in these amorphous ®lms [10], which is known to increase the optical energy gap E0 . Therefore, studying the effect of thermal annealing on the optical gap provides a better understanding of the mechanism of disorder and defect formation in the chalcogenide ®lms. The purpose of the present work is to study the structure and optical absorption of thin ®lms in the Ge20 Se80 x Bix system (0 x 10 at %). The compositional dependence of the optical energy gap of the investigated ®lms was studied. The effects of thermal
Permanent address: Department of Physics, Faculty of Science, Alexandria University, Egypt. Permanent address: Department of Physics, Faculty of Science, Assuit University, Egypt.
0957±4522
# 2001 Kluwer Academic Publishers
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annealing (below the crystallization temperature) on the optical properties as well as the structure of the ®lms were also examined.
2. Experimental
High purity (99.999%) selenium, germanium and bismuth (from Aldrich, USA) were used to prepare the chalcogenide glassy alloys. Bulk glasses of the Ge20 Se80 x Bix system (x 0, 2.5, 5, 7.5 and 10 at %) were prepared by the well established melt-quench technique. Thin ®lms were obtained by electron beam evaporation of the glassy bulk materials onto ultrasonically cleaned glass substrates, using an Edwards E-306 coating system operated at a base pressure of 6:7610 4 Pa. The substrate temperature was kept at room temperature during the evaporation process. The evaporation rate as well as the thickness of the investigated ®lms were controled using a quartz crystal monitor (Edward FTM5). For all the investigated compositions, a constant evaporation rate (2:5 nms 1 ) was used to deposit the required ®lms. An energy dispersive X-ray spectrometer (Link Analytical EDX) was used to measure the elemental composition of the investigated ®lms. The X-ray diffraction analysis was carried out for all as-deposited and thermally annealed ®lms using a Philips 1710 diffractometer (using CuKa radiation). The morphology and structure of the ®lms were determined using a Jeol 2000 transmission electron microscope (TEM) operated at 100 kV. Freshly evaporated as well as thermally treated ®lms were ¯oated off the substrates pre-coated with NaCl and carbon ®lms. They were ¯oated off by immersion in distilled water, then transferred to Cu microscope grids. Thermal treatment for TEM experiments was performed under a constant ¯ow of nitrogen gas. Optical re¯ectance
R and transmittance
T of the as-deposited and thermally annealed ®lms were measured at normal incidence at room temperature using a double-beam UV-VIS scanning spectrophotometer (SHIMADZU UV-2101) combined with a PC in the wavelength range 300±900 nm.
Figure 1 Glass forming region in the system Ge±Se±Bi; 2 2 after Tohge et al. [11]; after Pazin and Borisova [12]; s Composition of the asprepared Ge±Se±Bi glasses for the present work.
Bi content while the value of Tc decreases with increasing Bi content. The thermograms obtained for x 0 and 2.5 have lower values of Tg and do not show crystallization peaks. The X-ray diffraction analysis of all as-deposited thin ®lms indicated that they were in the amorphous state. The X-ray diffraction pattern of Ge20 Se70 Bi10 thin ®lm annealed at 593 K for one hour indicates two peaks at y 28.77 and 43.29, as shown in Fig. 3. The composition of the as-deposited Ge20 Se80 x Bix
3. Results
The phase diagram of the Ge-Se-Bi system [11, 12] is shown in Fig. 1. The phase diagram indicates that the largest amount of bismuth can be incorporated in the GeSe matrix when germanium has the content of 20 at %. Therefore, it was decided to investigate the series Ge20 Se80 x Bix (0 x 10 at %). It is clear that the investigated compositions lie within the glass-forming region. Differential scanning calorimetry (DSC) measurements were carried out for powder specimens of all compositions using a Du Pont 1090 thermal analyzer at a heating rate of 10 K min 1 . A typical DSC thermogram is shown in Fig. 2 for the composition Ge20 Se70 Bi10 . The thermogram indicates a glass transition temperature Tg at 468 K and two crystallization peaks Tc1 and Tc2 at 593 and 643 K. Similar thermograms were obtained for x 5 and 7.5, where the value of Tg increases with increasing 396
Figure 2 DSC thermogram for Ge20 Se70 Bi10 powder specimen (the heating rate is 10 K min 1 ).
thin ®lms (0 x 10 at %) was investigated using energy dispersive X-ray analysis (EDAX). Fig. 4 shows the spectral distribution of the constituent elements for the Ge20 Se70 Bi10 ®lm. The atomic percentage ratio of Ge, Se, and Bi were found to be 21.0, 68.9 and 10.1, respectively. The examination of the other compositions indicated a decrease of 1 at % in Se and an increase of 1 at % in Ge, while the Bi content is nearly the same as in the bulk starting material. The structure of Ge20 Se70 Bi10 thin ®lms was studied using transmission electron microscopy (TEM). The morphology as well as electron diffraction of the asdeposited and annealed Ge20 Se70 Bi10 ®lms of thickness 30 nm are shown in Fig. 5. The morphology of the asdeposited ®lms of different compositions was found to be of uniform contrast. Fig. 5a shows the electron diffraction pattern of the as-deposited Ge20 Se70 Bi10 ®lm. The diffused rings observed in the electron diffraction pattern indicate an amorphous structure. An amorphous structure is expected for the as-deposited Ge20 Se80 x Bix ®lms (0 x 10 at %) since the quenching rate during the deposition process is much higher than that of the investigated bulk alloys, as shown in the phase diagram (Fig. 1). TEM observations indicated amorphous-crystalline transformations for Ge20 Se70 Bi10 ®lms thermally annealed in situ at 593 K for 15 min. Homogeneously distributed interconnected polycrystalline phase is observed as a result of thermal annealing. The polycrystalline phase covered 60% of the ®lm surface as shown in the electron micrograph, Fig. 5b. The selected area electron diffraction (SAED) pattern for the transformed polycrystalline phase is shown in Fig. 5c and is characterized by diffraction rings. The analysis of the SAED pattern determines the most probable crystalline phases accompanying the transformations. The calibrated camera constant, obtained from the diffraction rings of the standard element gold, was used to calculate the d-spacings corresponding to different radii of the diffraction patterns of the annealed ®lms. The calculated d-spacings in the case of thermally annealed Ge20 Se70 Bi10 ®lm at 593 K are in fairly good agreement with the d-spacings of Bi2 Se3 [13]. In situ thermal annealing of the Ge20 Se70 Bi10 ®lm at 653 K showed the growth of dendritic phase as shown in the electron micrograph, Fig. 5d. (The micrograph was taken in situ during the annealing process.) Fig. 5e shows the amorphous-crystalline transformations after an incubation period of 15 min at 653 K. The new phase covers
Figure 3 X-ray diffraction pattern of Ge20 Se70 Bi10 thin ®lm thermally annealed at 593 K for 1 h.
Figure 4 Energy dispersive spectral distribution of the constituent elements for the as-deposited Ge20 Se70 Bi10 thin ®lm.
most of the specimen surface and grew with increasing annealing time at the same temperature. Fig. 5f shows the SAED for the transformed phase shown in Fig. 5e. The analysis of the SAED indicates the existence of Bi2 Se3 and BiSe crystalline phases [13, 14]. The spectral dependence of the optical transmittance and re¯ectance of the investigated Ge20 Se80 x Bix thin ®lms were obtained from the computerized double beam spectrophotometer in the wavelength range 300±900 nm. From the experimentally measured values of R and T, the absorption coef®cient
a was computed [15] according to the following expression: T
1
R2 exp
ad=1
R2 exp
2ad
1
where d is the thickness of the investigated ®lm (in cm). The above equation neglects the substrate effect and it can be applied only for homogeneous thin ®lms. In amorphous semiconductors, the optical absorption spectrum has been found to have three distinct regions [16]; namely, the high absorption region (a 104 cm 1 ), the exponential edge region (1 a 104 cm 1 ), and the weak absorption tail (a 1 cm 1 ) which originates from the defects and impurities. In the high absorption region, from which the optical band gap can be determined, the absorption is characterized by ahn B
hn
E0
2
2
where h is Planck's constant, n is the frequency, B is a parameter that depends on the transition probability and E0 is the optical energy gap of the investigated ®lm. The spectral variation at room temperature for the 1=2 absorption coef®cient plotted as
ahn versus the photon energy hn for the asdeposited and thermally annealed Ge20 Se75 Bi5 thin ®lms at 453 K fordifferent time intervals, is shown in Fig. 6. The linear relation of the
ahn1=2 versus hn plot indicates that the absorption mechanism in the Ge20 Se75 Bi5 thin ®lms is a non-direct transition [16]. The optical energy gap of the non-direct transition can be obtained from the intercept of the 1=2
ahn versus hn plots with the energy axis at
ahn1=2 0. Fig. 7 represents the variation of the optical gap E0 for Ge20 Se75 Bi5 thin ®lms thermally annealed at 453 K as a function of annealing time. The optical energy gap increases gradually from 1.2 to 1.33 eV with increasing 397
(a)
(b)
(c)
(d)
(e)
(f )
Figure 5 (a) Electron diffraction pattern for the as-deposited Ge20 Se70 Bi10 thin ®lm, (b) TEM micrograph of Ge20 Se70 Bi10 thin ®lm thermally annealed in situ at 593 K for 15 min, (c) SAED pattern for Ge20 Se70 Bi10 thin ®lm thermally annealed in situ at 593 K for 15 min, (d) TEM micrograph showing the growth of dendritic phase appeared on Ge20 Se70 Bi10 thin ®lm after annealing at 653 K, (e) TEM micrograph showing the crystallization of Ge20 Se70 Bi10 thin ®lm thermally annealed in situ at 653 K for 15 min, (f ) SAED pattern for the crystallized phase shown in Fig. 5e.
annealing time, then it reaches a steady value. Our experimental measurements indicate that the optical energy gap E0 increases gradually with thermal annealing in the temperature range 373 to 453 K before it attains the steady value. The time needed for steady optical gap decreases with increasing annealing temperature. Films of different compositions were thermally annealed at 413 K for one hour which is below the glass transition temperature before optical measurements at room temperature were carried out. The spectral 398
variation at room temperature for the absorption coef®cient plotted as
ahn1=2 versus the photon energy hn for Ge-Se-Bi thin ®lms of different Bi content is shown in Fig. 8. It is clear that the absorption coef®cient moves towards lower values of hn with increasing Bi content. The calculated values of the optical energy gap were plotted versus Bi content (in at %) as shown in Fig. 9. The addition of 2.5 at % of Bi to Ge20 Se80 decreases E0 from 1.93 to 1.55 eV. The rate of the change of the optical gap decreases as the Bi content increases.
Figure 6
ahn1=2 as a function of photon energy hn for Ge20 Se75 Bi5 thin ®lm heat treated at 453 K for different periods of time.
Figure 8 The absorption coef®cient plotted as
ahn1=2 as a function of photon energy hn for Ge20 Se80 x Bix thin ®lms.
The plots given in Figs 6 and 8 indicate that the absorption in Ge20 Se80 x Bix ®lms (0 x 10 at %) is due to non-direct transitions and the extrapolation of the linear portion of the curves to the abscissa has been used to ®nd the optical energy gap E0 . The composition dependence of the optical energy gap (Fig. 9) indicates that the values of E0 decrease continuously with increasing Bi content. The values of the optical energy gap almost agrees with that obtained by other investigators for the bulk glasses [11] and ¯ash-evaporated ®lms [17], and might be explained on the basis of increased band-tailing [18]. This behavior could be understood from the standpoint of chemical bonds formed in these
4. Discussion
glasses. The following assumptions [19] were used for the chemical bond approach: (1) Se bonds divalently and Ge bonds tetravalently, which implies that the dangling bonds and other valence defects have been neglected to a ®rst approximation. Also, van der Waals interactions are neglected. On the other hand, the number of covalent bonds per Bi atom is assumed to be six, since Bi is known to be 6-fold coordinated by chalcogen in Bi2 Se3 and Bi2 Te3 crystals [19, 20]. (2) Atoms would combine more favorably with atoms of different kinds than with the same kind. This means that the maximum amount of chemical ordering has been assumed. Bonds between like atoms will occur only if there is an excess of a certain type of atom, so that it is not possible to satisfy its
Figure 7 Optical energy gap E0 as a function of annealing time for Ge20 Se75 Bi5 thin ®lm annealed at 453 K.
Figure 9 The dependence of the optical gap on Bi content (in at %) for Ge20 Se80 x Bix thin ®lms.
399
T A B L E I The bond energy calculated for different bonds (A±B)
D(A±B) (kcal mol
Ge±Se Se±Bi Se±Se Bi±Ge Ge±Ge Bi±Bi
49.4 40.7 44.0 31.0 37.6 25.0
1
)
valence requirements by bonding to other atoms of different kinds. (3) Bonds are formed in the sequence of decreasing bond energy. The bond energy D(A±B) for a heteronuclear bond was calculated following the relation which was proposed by Pauling [21]: D
A--B D
A--A ? D
B--B1=2 30
xA
xB 2
3
where D(A±A) and D(B±B) are the energies of the homonuclear bonds and xA ; xB are the electronegativities of the atoms involved. The D(A±A) values (in units of kcal/mol) for Ge, Se and Bi are 37.6, 44 and 25, respectively. The xB values for Ge, Se and Bi are 1.8, 2.4 and 1.8, respectively [21]. The bond energies, calculated from Equation 3 for all bonds, are given in Table I. When the chemical bond approach is applied to the Ge20 Se80 specimen, it is expected that only the Ge±Se and Se±Se bonds would be present. When Bi is incorporated into the matrix, Bi is expected to bond preferably with Se and not to Ge. As for the Ge20 Se75 Bi5 specimen, the following picture would be expected: Since the Ge±Se bond has the largest bond energy value, therefore it is expected that Se atoms would ®rst completely saturate the 80 valences of the 20 Ge atoms and then completely saturate the 30 valences of the Bi atoms. Then the excess Se atoms would have no other choice but to satisfy their remaining valence requirements by bonding among themselves (40 valences). Therefore the Ge±Bi, Ge±Ge and Bi±Bi bonds are not expected to form in this composition. The chemical bonds approach predicts that only three types of bonds, namely Ge±Se, Se±Bi and Se±Se bonds, are expected to form for the investigated compositions. In other words, the Ge±Bi and Ge±Ge bonds are not present. This is in agreement with XPS measurements [22], where there was no evidence of any Ge±Ge or Ge± Bi bonds in the investigated Ge±Se±Bi ®lms. Also the measurements of the X-ray GeK absorption study of Ge20 Se80 x Bix (x 0, 7 and 13 at %) [23], support the absence of Ge±Bi bonds. On the other hand, Table II indicates that the number of Ge±Se bonds is the same for all compositions. This is
Figure 10 Density of states g
E as a function of energy E. Shaded regions are localized states.
in agreement with the measurements of Raman spectra of sputtered amorphous Ge25 Se75 x Bix ®lms (0 x 19 at %) [24], where it was found that the incorporation of Bi does not affect the Ge±Se bonding. The decrease of the optical energy gap E0 with the increase of Bi content is analogous to that of the activation energy of d.c. conduction [11], indicating the increase of the density of localized states with increasing Se content. The results shown in Figs 6 and 7 indicate that thermal annealing in the temperature range from 353±473 K causes an increase in the nondirect optical energy gap for all compositions. A similar increase in E0 with annealing has been observed in Ge±Se glasses [25] and also for other glasses [26]. Such a dependence on annealing is attributed to the reduction in disorder in the atomic bonding between neighbors and hence the decrease of the density of tail states adjacent to the band gap. The behavior of the optical gap with annealing temperature in amorphous Ge20 Se80 x Bix ®lms (0 5 x 10 at %) could be explained by applying the model of the density of states in amorphous solids, which was proposed by Mott and Davis [27]. Considering Fig. 10, the width of localized states near the mobility edges
EC EA and
EB EV depend on the degree of disorder or on defects in the amorphous structure. In particular, it is known that unsaturated bonds are responsible for the formation of some defects in the
T A B L E I I The number of Ge±Se, Se±Bi and Se±Se bonds calculated for all compositions Composition
Number of Ge±Se bonds
Number of Se±Bi bonds
Number of Se±Se bonds
Ge20 Se80 Ge20 Se77:5 Bi2:5 Ge20 Se75 Bi5 Ge20 Se72:5 Bi7:5 Ge20 Se70 Bi10
80 80 80 80 80
0 15 30 45 60
80 60 40 20 0
400
amorphous solids. Such defects produce localized states in the band structure. The presence of a high concentration of localized states is responsible for the relatively low values of the optical gap in the case of asprepared amorphous ®lms. During the heat-treatment process, the unsaturated defects are gradually annealed out, thus producing a larger number of saturated bonds. This decreases the density of localized states, and therefore increases the optical gap.
5. Conclusions
Structural and optical investigations on Ge20 Se80 x Bix thin ®lms (0 x 10 at %) were carried out. Thermal annealing (at temperatures below the crystallization temperature) causes an increase in the non-direct optical energy gap for all compositions. This is attributed to the decrease of the density of tail states adjacent to the band gap. The morphology of the as-prepared ®lms was found to be of uniform contrast. The electron diffraction patterns were characterized by diffused rings, indicative of the amorphous state of the ®lms. Thermal annealing at higher temperatures leads to crystallization as indicated by the TEM observations as well as the electron diffraction patterns, which was indicated by the diffraction rings. It was found that the absorption in the Ge20 Se80 x Bix thin ®lms (0 x 10 at %) is due to non-direct transitions. The optical gap E0 was found to decrease with increasing Bi content. This trend was analyzed by using the chemical bonds approach. According to this model, it was suggested that the addition of Bi leads to the formation of Se±Bi bonds at the expense of the Se±Se bonds. The number of the Se±Bi bonds increases with increasing Bi content.
3. 4. 5. 6. 7. 8.
9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25.
Acknowledgments
The authors are thankful to the ``Research Council Funding'', The United Arab Emirates University for the ®nancial support (Grant number 06-2-11/99).
26. 27.
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Received 21 August 2000 and accepted 15 March 2001
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