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Controlling the growth of ultrasmall CdTe quantum dots and the diffusion of cadmium vacancies: Thermal annealing
Journal of Alloys and Compounds 637 (2015) 466–470
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
Controlling the growth of ultrasmall CdTe quantum dots and the diffusion of cadmium vacancies: Thermal annealing Noelio O. Dantas ⇑, Guilherme L. Fernandes, Anielle Christine A. Silva ⇑ Laboratório de Novos Materiais Isolantes e Semicondutores (LNMIS), Instituto de Física, Universidade Federal de Uberlândia, Uberlândia, Brazil
a r t i c l e
i n f o
Article history: Received 20 November 2014 Received in revised form 21 February 2015 Accepted 4 March 2015 Available online 12 March 2015 Keywords: Ultrasmall CdTe quantum dots Silicate glass matrix Diffusion of cadmium vacancies Optical properties
a b s t r a c t Ultrasmall CdTe quantum dots (USQDs) were successfully grown in a silicate glass matrix by fusion and after thermal annealing. Growth control of USQDs was investigated by optical absorption (OA), atomic force microscopy (AFM), transmission electron microscopy (TEM) and photoluminescence (PL). A redshift of OA band with increasing thermal annealing time provided evidence of CdTe USQD growth. This increase of average size of the CdTe USQDs was determined by OA spectra, AFM and TEM images. In addition, PL spectra showed that longer thermal annealing times decreased deep levels luminescent intensity from cadmium vacancies (VCd) in the CdTe USQDs. This phenomenon occurred because VCd diffused to the USQDs’ surface with longer thermal annealing times. Therefore, we control the growth of CdTe USQDs as well as the luminescent intensity from surface defects and VCd as a function of thermal annealing time. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Cadmium telluride (CdTe) is a II–VI semiconductor with a zincblend crystalline structure, exciton Bohr radius of 6.5 nm, absorbs and emits in the infrared optical window with band gap energy of 1.5 eV (826.5 nm) at room temperature [1] has a high melting point of 1041 °C with evaporation starting at 1050 °C. Usually, CdTe shows the p-type semiconductor because the cadmium vacancies (VCd) are present [2]. One of the exceptional advantages of semiconductor nanocrystals (NCs) is that physical properties are changed by the size and shape. Thus, promising materials simply tuning the absorption and luminescence spectra of the NCs aiming several optical applications [3]. CdTe quantum dots (QDs) are often used as luminescent probes [4–7], solar cells [8,9], photodetectors [10,11] and photovoltaic devices [12,13]. Depending on the application, QDs can be inserted into various systems such as aqueous solutions [14,15], polymer films [16–20], substrates [21–23] and glass matrix [24– 28]. QDs synthesized in glass matrix has several advantages, such as high mechanical and chemical stability, protection of the semiconductor material against the environment, prevention of agglomeration, and adaptability to device manufacturing processes [29,30]. These characteristics allow QDs to have potential applications in optoelectronics and optical cut-off filters [10,30–32].
The lower the quantum dots more the quantum confinement properties are enhanced. But it is noteworthy that this size is related to the Bohr radius of each material. Ultrasmall CdTe QDs are still little studied and mostly obtained by solutions methods [33–35]. Thus, in this study, probably for the first time, we controlled the growth of ultrasmall CdTe QDs (USQDs) in a silicate glass matrix synthesized by fusion as a function of thermal annealing time. The optical and morphological properties of the CdTe USQDs embedded in the silicate glass matrix were investigated by optical absorption (OA), atomic force microscopy (AFM), scanning transmission electron microscopy (TEM) and photoluminescence (PL). 2. Experimental 2.1. Materials CdTe USQDs were synthesized in SNAB glass matrix with a nominal composition of 40SiO2 30Na2CO3 1Al2O3 29B2O3 (mol%) and 4CdTe bulk (wt%). Preparation consisted of melting the powder mixtures in an alumina crucible at 1250 °C for 15 min. The melt was then quickly cooled to 0 °C. Next, thermal annealing was carried out at 555 °C for 0, 2 and 10 h, since, normally we adopted around 30 °C above the glass transition temperature [Tg (SNAB) = 525.15 °C = 798.15 K], which is the minimum energy required to cause molecular mobility and the diffusion of precursor ions (Cd2+ and Te2) for the formation/growth of nanocrystals (CdTe USQDs) [35]. 2.2. Characterization
⇑ Corresponding authors. E-mail addresses: [email protected] (N.O. Dantas), [email protected] (A.C.A. Silva). http://dx.doi.org/10.1016/j.jallcom.2015.03.026 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
Optical absorption (OA) spectra were measured with a UV-3600 UV–VIS–NIR Shimadzu spectrometer operating from 190 to 3300 nm and with a spectral resolution of 1 nm. Atomic force microscopy (AFM) images of the USQD samples were
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recorded at room temperature using a scanning probe microscope (Shimadzu, SPM–9600). The AFM technique was also used to estimate the average size of the as-synthesized QDs based on height distributions. It is important to mention that the SPM-9600 microscope is precise to 0.01 nm in the z-direction (height) [36]. First, the glass sample was polished meticulously and thoroughly to ensure that irregularities in the glass surface were smaller than the average USQD size. Polishing was necessary to obtain AFM images showing a regular distribution of USQDs on the flat surface of the glass matrix. In addition, the glass sample was cleaned immediately before the AFM measurement to eliminate contaminants on the polished surface, which could reduce image quality. Scanning transmission electron micrographs (TEM, JEOL, JEM-2100, 200 kV) were taken to investigate the formation and size of the CdTe USQDs. Photoluminescence (PL) spectra were recorded with Jaz spectrometer (Ocean Optics Inc.), using a excitation wavelength of 409 nm from diode laser.
3. Results and discussion Fig. 1 shows the (a) photographs and (b) OA spectra of the SNAB glass matrix and samples doped with CdTe post thermal annealing at 555 °C for 0, 2 and 10 h. In the photographs observe that SNAB matrix is transparent, already the sample doped with CdTe is slightly dark, increasing thermal annealing time the samples are darker. This result gives indications that the increase in thermal annealing time favors the growth of the nanocrystal. The OA spectrum of the SNAB glass matrix is transparent in the visible range where CdTe absorbs and emits (Fig. 1(b)). Thus, it is essential to optically observe the control and growth of CdTe NCs. In the sample doped with CdTe with and without thermal treatment observed an OA band in the spectra. This band is at higher energy than that of bulk CdTe (1.44 eV–860 nm), which indicates that the NCs exhibit quantum confinement effects called quantum dots [3]. The OA band from the sample without thermal annealing confirms that the cooling rate, which the melt was submitted, not as fast enough to prevent the diffusion of ions of Cd2+ and Te2 and the formation of CdTe USDQs during cooling/formation glass. The small redshift in the OA band with increasing thermal annealing time is a sign of quantum confinement in dots of increasing size. In addition, the sample that was thermally annealed for 10 h shows a broadening of the OA band, which indicates greater size dispersion. The increase of size dispersion is due to growth competition among different sized NCs within the SNAB glass matrix [37]. The maximum values of the OA bands were acquired from Gaussian fit. The average size of the CdTe QDs was determined by effective mass approximation (EMA) where transition energy between electron and 2
2
2
heavy-hole states was estimated using EConf ¼ Egap þ 2hlpR2 1:8 eeR where Eg (EgCdTe = 1.475 eV) [38] is the bulk semiconductor energy gap, l the electron–hole reduced effective mass (lCdTe = 0.068), [38] and R is the radius of the spherical confinement region. The third term is an estimate of the electron–hole Coulomb interaction where e (eCdTe = 10.4) [39] is the dielectric constant. Estimated average sizes (R) for the CdTe QDs were, 2.01 nm, 2.04 nm and 2.06 nm for thermal annealing times of 0, 2 and 10 h, respectively. These values are within the range of ultrasmall CdTe QDs (USQDs) and are in excellent agreement with those obtained by the AFM and TEM images (see Figs. 2 and 3) [40–43]. Thus, probably, for the first time, the growth of CdTe USQDs was controlled within a silicate glass matrix and the size of these USQDs was manipulated by simply changing thermal annealing time. The formation of CdTe USQDs through AFM images is shows in Fig. 2, where (a) SNAB matrix and samples doped with CdTe and then thermally annealed at 555 °C for (b) 0 and (c) 10 h. The AFM images of the SNAB matrix do not show crystal formation; however, the images of the samples doped with CdTe show NC formation (Fig. 2b and c). These results are in excellent agreement with the OA spectra and confirm the formation of USQDs (Fig. 1) [40]. The AFM images show that the CdTe USQDs are spherically shaped and mostly homogeneous, despite occasional deviations.
Fig. 1. (a) Photographs, (b) room temperature OA spectra of SNAB glass matrix and samples doped with CdTe post thermal annealing at 555 °C for 0, 2 and 10 h.
The average radius (R) of the CdTe USQDs was obtained from the height distribution of the AFM images (R = 2.01 nm and R = 2.11 nm to samples thermally annealed at 555 °C for 0 and 10 h, respectively), as showed in Fig. 2 [36]. Furthermore, it is well-established that the mean USQD size estimated from topographic images is equal to that of USQDs embedded in our sample [36]. The average radius values obtained are in agreement with those calculated from the OA spectra and based on effective mass approximation. The size of the CdTe USQDs increases as thermal annealing time (histograms in Fig. 2b and c). This observation is in agreement with the small redshift in the OA band as annealing time increased (see Fig. 1). Therefore, longer thermal annealing times affect nucleation and produce larger CdTe USQDs. In order to confirm the size of the NCs we performed TEM analysis as show in Fig. 3. The TEM images showed NCs spherical, uniformly sized, and monodispersed in the glass matrix, beside the effect of thermal annealing time. The increase of the thermal annealing time produces larger QDs (4.2–4.6 nm for 0 h and 10 h,
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(753 nm–1.65 eV) and E2 (892 nm–1.39 eV) are due intrinsic lattice defects in CdTe NCs namely cadmium vacancies (VCd) on different crystallographic position [46]. In the PL spectra observed that the luminescence these bands decreased with longer thermal annealing time. It is occurred due to the diffusion of VCd to the surface of CdTe USQDs with increasing thermal annealing time that is supported by the fact that the NCs is in thermodynamic equilibrium with its environment, and chemical equilibrium with reservoirs of impurities and NCs constituents [48–50]. This diffusion of VCd with increased of the thermal annealing time was also confirmed in our recent paper in CdSe QDs [48]. Therefore, the impure atoms must be able to diffuse readily through the NCs to become purified. This process is called ‘‘self-purification’’ and occurs for different types of impurities such as defects and ions [49,51]. The diagram in Fig. 4(b) depicts the emissions detected in the PL spectra and identifies the four emission bands. Fig. 4(c) shows the ratio of the deep level to the excitonic level emission intensities. The E1/Eexn and E2/ Eexn ratios decrease with increasing thermal annealing times. Additionally, the E1/Eexn ratio decreases more sharply than the E2/ Eexn ratio, showing that the VCds corresponding to the E1 emission diffuse more easily to the surface of the CdTe USQDs than do the VCds corresponding to the E2 emission. Thus, is difference between E1 and E2 is due to the strong dependence of the crystallographic position in which the vacancy is located [52,53]. Therefore, we controlled VCd diffusion by controlling thermal annealing time and consequently reduced the deep level densities corresponding E1 and E2 emissions. Fig. 4(d) shows the ratio of the surface defect to the excitonic level emission intensities (ESDL/Eexn). The ratio ESDL/Eexn decrease with increasing thermal annealing times. This result is consistent with increase of size NCs, in excellent agreement with Figs. 1–3. This decrease was expected because formation of surface defects in NCs decreases as size increases [47,54]. Therefore, on these results, we can say that controlled the formation and growth
Fig. 2. AFM images of (a) SNAB glass matrix and samples doped with CdTe and treated with post thermal annealing at 555 °C for (b) 0 and (c) 10 h.
respectively). This result is in excellent agreement with the values obtained based on the EMA from the spectra AO and those of the histograms from the AFM images. In the inset observed that the average d-spacing [1 1 1] of the NCs is around 0.374 nm, in accordance with the cubic-structured CdTe that confirm the grown CdTe USQDs [44]. The normalized PL spectra of the SNAB glass matrix and samples doped with CdTe post thermal annealing at 555 °C for 0, 2 and 10 h are shows Fig. 4(a). All PL spectra of the samples doped with CdTe show four PL bands at 529 nm (2.34 eV), 564 nm (2.19 eV), 753 nm (1.65 eV) and 892 nm (1.39 eV), which correspond to the excitonic (Eexn) of surface defect (ESDL) and deep level (E1 and E2) emissions, respectively [45,46]. It is interesting to note that the Eexn bands in the PL spectra are shifted slightly to longer wavelengths with longer thermal annealing times. This result is in agreement with the OA spectra (Fig. 1), AFM (Fig. 2) and TEM (Fig. 3) images, confirming that longer annealing times are associated with increase of QDs size. Thus, the positions of the Eexn bands are strongly dependent on thermal annealing times. The surface defect level (ESDL = 564 nm (2.19 eV)) are from the lattice defects near surface of the NCs, which are due to their much greater surface to volume ratio and the presence of interface [47]. The deep level emissions labeled E1
Fig. 3. TEM images of SNAB matrix and samples doped with CdTe and treated with post thermal annealing at 555 °C for (a) 0 and (b) 10 h. These images illustrate the morphology, average diameter size and average d-spacing [1 1 1] of the CdTe USQDs.
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Acknowledgements The authors gratefully acknowledge financial support from the agencies CAPES, FAPEMIG and MCT/CNPq. AFM Shimadzu at the Institute of Physics (INFIS) supported by the grant ‘‘PróEquipamentos’’ from the Brazilian Agency CAPES. References
Fig. 4. PL spectra of (a) the SNAB glass matrix and samples doped with CdTe with post thermal annealing at 555 °C for 0, 2 and 10 h. (b) Different recombination processes, where the emissions from CdTe USQDs are identified. The Eexn (2.34 eV), ESDL (2.19 eV), E1 (1.65 eV) and E2 (1.39 eV) emissions refer to the excitonic emission and defect levels present in the CdTe USQDs. The non-radiative processes are also indicated. (c) Ratio of the defect to the excitonic level emission intensities, (ESDL/Eexn), E1/Eexn and E2/Eexn with increasing thermal annealing times. (d) Ratio of the surface defect to the excitonic level emission intensities (ESDL/Eexn) with increasing thermal annealing times.
of CdTe USQDs and the luminescent intensities arising from surface defects and VCd due to the heat treatment time. 4. Conclusions We grew CdTe USQDs in a SNAB (40SiO2 30Na2CO3 1Al2O3 29B2O3 (mol%)) glass matrix by the fusion method and controlled the growth of the CdTe USQDs by varying thermal annealing time. Longer thermal annealing times caused refd shift of the PL and OA bands, indicating growth of USQDs. The average size of the USQDs in the AFM and TEM images confirmed this phenomenon and was in agreement with the size estimated by the EMA model. We also observed that longer thermal annealing time favored decreased luminescence corresponding to the deep level of the VCd within the CdTe USQDs. This was caused in turn by the diffusion of VCd to the surface of CdTe USQDs. Moreover, the decrease of the surface defect level luminescence was observed confirming the increases of size NCs. Therefore, we also controlled VCd diffusion as a function of thermal annealing time. We hope these results stimulate further research into the growth of several types USQDs with desirable properties for optoelectronic device applications.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
Controlling the growth of ultrasmall CdTe quantum dots and the diffusion of cadmium vacancies: Thermal annealing Noelio O. Dantas ⇑, Guilherme L. Fernandes, Anielle Christine A. Silva ⇑ Laboratório de Novos Materiais Isolantes e Semicondutores (LNMIS), Instituto de Física, Universidade Federal de Uberlândia, Uberlândia, Brazil
a r t i c l e
i n f o
Article history: Received 20 November 2014 Received in revised form 21 February 2015 Accepted 4 March 2015 Available online 12 March 2015 Keywords: Ultrasmall CdTe quantum dots Silicate glass matrix Diffusion of cadmium vacancies Optical properties
a b s t r a c t Ultrasmall CdTe quantum dots (USQDs) were successfully grown in a silicate glass matrix by fusion and after thermal annealing. Growth control of USQDs was investigated by optical absorption (OA), atomic force microscopy (AFM), transmission electron microscopy (TEM) and photoluminescence (PL). A redshift of OA band with increasing thermal annealing time provided evidence of CdTe USQD growth. This increase of average size of the CdTe USQDs was determined by OA spectra, AFM and TEM images. In addition, PL spectra showed that longer thermal annealing times decreased deep levels luminescent intensity from cadmium vacancies (VCd) in the CdTe USQDs. This phenomenon occurred because VCd diffused to the USQDs’ surface with longer thermal annealing times. Therefore, we control the growth of CdTe USQDs as well as the luminescent intensity from surface defects and VCd as a function of thermal annealing time. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Cadmium telluride (CdTe) is a II–VI semiconductor with a zincblend crystalline structure, exciton Bohr radius of 6.5 nm, absorbs and emits in the infrared optical window with band gap energy of 1.5 eV (826.5 nm) at room temperature [1] has a high melting point of 1041 °C with evaporation starting at 1050 °C. Usually, CdTe shows the p-type semiconductor because the cadmium vacancies (VCd) are present [2]. One of the exceptional advantages of semiconductor nanocrystals (NCs) is that physical properties are changed by the size and shape. Thus, promising materials simply tuning the absorption and luminescence spectra of the NCs aiming several optical applications [3]. CdTe quantum dots (QDs) are often used as luminescent probes [4–7], solar cells [8,9], photodetectors [10,11] and photovoltaic devices [12,13]. Depending on the application, QDs can be inserted into various systems such as aqueous solutions [14,15], polymer films [16–20], substrates [21–23] and glass matrix [24– 28]. QDs synthesized in glass matrix has several advantages, such as high mechanical and chemical stability, protection of the semiconductor material against the environment, prevention of agglomeration, and adaptability to device manufacturing processes [29,30]. These characteristics allow QDs to have potential applications in optoelectronics and optical cut-off filters [10,30–32].
The lower the quantum dots more the quantum confinement properties are enhanced. But it is noteworthy that this size is related to the Bohr radius of each material. Ultrasmall CdTe QDs are still little studied and mostly obtained by solutions methods [33–35]. Thus, in this study, probably for the first time, we controlled the growth of ultrasmall CdTe QDs (USQDs) in a silicate glass matrix synthesized by fusion as a function of thermal annealing time. The optical and morphological properties of the CdTe USQDs embedded in the silicate glass matrix were investigated by optical absorption (OA), atomic force microscopy (AFM), scanning transmission electron microscopy (TEM) and photoluminescence (PL). 2. Experimental 2.1. Materials CdTe USQDs were synthesized in SNAB glass matrix with a nominal composition of 40SiO2 30Na2CO3 1Al2O3 29B2O3 (mol%) and 4CdTe bulk (wt%). Preparation consisted of melting the powder mixtures in an alumina crucible at 1250 °C for 15 min. The melt was then quickly cooled to 0 °C. Next, thermal annealing was carried out at 555 °C for 0, 2 and 10 h, since, normally we adopted around 30 °C above the glass transition temperature [Tg (SNAB) = 525.15 °C = 798.15 K], which is the minimum energy required to cause molecular mobility and the diffusion of precursor ions (Cd2+ and Te2) for the formation/growth of nanocrystals (CdTe USQDs) [35]. 2.2. Characterization
⇑ Corresponding authors. E-mail addresses: [email protected] (N.O. Dantas), [email protected] (A.C.A. Silva). http://dx.doi.org/10.1016/j.jallcom.2015.03.026 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
Optical absorption (OA) spectra were measured with a UV-3600 UV–VIS–NIR Shimadzu spectrometer operating from 190 to 3300 nm and with a spectral resolution of 1 nm. Atomic force microscopy (AFM) images of the USQD samples were
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recorded at room temperature using a scanning probe microscope (Shimadzu, SPM–9600). The AFM technique was also used to estimate the average size of the as-synthesized QDs based on height distributions. It is important to mention that the SPM-9600 microscope is precise to 0.01 nm in the z-direction (height) [36]. First, the glass sample was polished meticulously and thoroughly to ensure that irregularities in the glass surface were smaller than the average USQD size. Polishing was necessary to obtain AFM images showing a regular distribution of USQDs on the flat surface of the glass matrix. In addition, the glass sample was cleaned immediately before the AFM measurement to eliminate contaminants on the polished surface, which could reduce image quality. Scanning transmission electron micrographs (TEM, JEOL, JEM-2100, 200 kV) were taken to investigate the formation and size of the CdTe USQDs. Photoluminescence (PL) spectra were recorded with Jaz spectrometer (Ocean Optics Inc.), using a excitation wavelength of 409 nm from diode laser.
3. Results and discussion Fig. 1 shows the (a) photographs and (b) OA spectra of the SNAB glass matrix and samples doped with CdTe post thermal annealing at 555 °C for 0, 2 and 10 h. In the photographs observe that SNAB matrix is transparent, already the sample doped with CdTe is slightly dark, increasing thermal annealing time the samples are darker. This result gives indications that the increase in thermal annealing time favors the growth of the nanocrystal. The OA spectrum of the SNAB glass matrix is transparent in the visible range where CdTe absorbs and emits (Fig. 1(b)). Thus, it is essential to optically observe the control and growth of CdTe NCs. In the sample doped with CdTe with and without thermal treatment observed an OA band in the spectra. This band is at higher energy than that of bulk CdTe (1.44 eV–860 nm), which indicates that the NCs exhibit quantum confinement effects called quantum dots [3]. The OA band from the sample without thermal annealing confirms that the cooling rate, which the melt was submitted, not as fast enough to prevent the diffusion of ions of Cd2+ and Te2 and the formation of CdTe USDQs during cooling/formation glass. The small redshift in the OA band with increasing thermal annealing time is a sign of quantum confinement in dots of increasing size. In addition, the sample that was thermally annealed for 10 h shows a broadening of the OA band, which indicates greater size dispersion. The increase of size dispersion is due to growth competition among different sized NCs within the SNAB glass matrix [37]. The maximum values of the OA bands were acquired from Gaussian fit. The average size of the CdTe QDs was determined by effective mass approximation (EMA) where transition energy between electron and 2
2
2
heavy-hole states was estimated using EConf ¼ Egap þ 2hlpR2 1:8 eeR where Eg (EgCdTe = 1.475 eV) [38] is the bulk semiconductor energy gap, l the electron–hole reduced effective mass (lCdTe = 0.068), [38] and R is the radius of the spherical confinement region. The third term is an estimate of the electron–hole Coulomb interaction where e (eCdTe = 10.4) [39] is the dielectric constant. Estimated average sizes (R) for the CdTe QDs were, 2.01 nm, 2.04 nm and 2.06 nm for thermal annealing times of 0, 2 and 10 h, respectively. These values are within the range of ultrasmall CdTe QDs (USQDs) and are in excellent agreement with those obtained by the AFM and TEM images (see Figs. 2 and 3) [40–43]. Thus, probably, for the first time, the growth of CdTe USQDs was controlled within a silicate glass matrix and the size of these USQDs was manipulated by simply changing thermal annealing time. The formation of CdTe USQDs through AFM images is shows in Fig. 2, where (a) SNAB matrix and samples doped with CdTe and then thermally annealed at 555 °C for (b) 0 and (c) 10 h. The AFM images of the SNAB matrix do not show crystal formation; however, the images of the samples doped with CdTe show NC formation (Fig. 2b and c). These results are in excellent agreement with the OA spectra and confirm the formation of USQDs (Fig. 1) [40]. The AFM images show that the CdTe USQDs are spherically shaped and mostly homogeneous, despite occasional deviations.
Fig. 1. (a) Photographs, (b) room temperature OA spectra of SNAB glass matrix and samples doped with CdTe post thermal annealing at 555 °C for 0, 2 and 10 h.
The average radius (R) of the CdTe USQDs was obtained from the height distribution of the AFM images (R = 2.01 nm and R = 2.11 nm to samples thermally annealed at 555 °C for 0 and 10 h, respectively), as showed in Fig. 2 [36]. Furthermore, it is well-established that the mean USQD size estimated from topographic images is equal to that of USQDs embedded in our sample [36]. The average radius values obtained are in agreement with those calculated from the OA spectra and based on effective mass approximation. The size of the CdTe USQDs increases as thermal annealing time (histograms in Fig. 2b and c). This observation is in agreement with the small redshift in the OA band as annealing time increased (see Fig. 1). Therefore, longer thermal annealing times affect nucleation and produce larger CdTe USQDs. In order to confirm the size of the NCs we performed TEM analysis as show in Fig. 3. The TEM images showed NCs spherical, uniformly sized, and monodispersed in the glass matrix, beside the effect of thermal annealing time. The increase of the thermal annealing time produces larger QDs (4.2–4.6 nm for 0 h and 10 h,
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(753 nm–1.65 eV) and E2 (892 nm–1.39 eV) are due intrinsic lattice defects in CdTe NCs namely cadmium vacancies (VCd) on different crystallographic position [46]. In the PL spectra observed that the luminescence these bands decreased with longer thermal annealing time. It is occurred due to the diffusion of VCd to the surface of CdTe USQDs with increasing thermal annealing time that is supported by the fact that the NCs is in thermodynamic equilibrium with its environment, and chemical equilibrium with reservoirs of impurities and NCs constituents [48–50]. This diffusion of VCd with increased of the thermal annealing time was also confirmed in our recent paper in CdSe QDs [48]. Therefore, the impure atoms must be able to diffuse readily through the NCs to become purified. This process is called ‘‘self-purification’’ and occurs for different types of impurities such as defects and ions [49,51]. The diagram in Fig. 4(b) depicts the emissions detected in the PL spectra and identifies the four emission bands. Fig. 4(c) shows the ratio of the deep level to the excitonic level emission intensities. The E1/Eexn and E2/ Eexn ratios decrease with increasing thermal annealing times. Additionally, the E1/Eexn ratio decreases more sharply than the E2/ Eexn ratio, showing that the VCds corresponding to the E1 emission diffuse more easily to the surface of the CdTe USQDs than do the VCds corresponding to the E2 emission. Thus, is difference between E1 and E2 is due to the strong dependence of the crystallographic position in which the vacancy is located [52,53]. Therefore, we controlled VCd diffusion by controlling thermal annealing time and consequently reduced the deep level densities corresponding E1 and E2 emissions. Fig. 4(d) shows the ratio of the surface defect to the excitonic level emission intensities (ESDL/Eexn). The ratio ESDL/Eexn decrease with increasing thermal annealing times. This result is consistent with increase of size NCs, in excellent agreement with Figs. 1–3. This decrease was expected because formation of surface defects in NCs decreases as size increases [47,54]. Therefore, on these results, we can say that controlled the formation and growth
Fig. 2. AFM images of (a) SNAB glass matrix and samples doped with CdTe and treated with post thermal annealing at 555 °C for (b) 0 and (c) 10 h.
respectively). This result is in excellent agreement with the values obtained based on the EMA from the spectra AO and those of the histograms from the AFM images. In the inset observed that the average d-spacing [1 1 1] of the NCs is around 0.374 nm, in accordance with the cubic-structured CdTe that confirm the grown CdTe USQDs [44]. The normalized PL spectra of the SNAB glass matrix and samples doped with CdTe post thermal annealing at 555 °C for 0, 2 and 10 h are shows Fig. 4(a). All PL spectra of the samples doped with CdTe show four PL bands at 529 nm (2.34 eV), 564 nm (2.19 eV), 753 nm (1.65 eV) and 892 nm (1.39 eV), which correspond to the excitonic (Eexn) of surface defect (ESDL) and deep level (E1 and E2) emissions, respectively [45,46]. It is interesting to note that the Eexn bands in the PL spectra are shifted slightly to longer wavelengths with longer thermal annealing times. This result is in agreement with the OA spectra (Fig. 1), AFM (Fig. 2) and TEM (Fig. 3) images, confirming that longer annealing times are associated with increase of QDs size. Thus, the positions of the Eexn bands are strongly dependent on thermal annealing times. The surface defect level (ESDL = 564 nm (2.19 eV)) are from the lattice defects near surface of the NCs, which are due to their much greater surface to volume ratio and the presence of interface [47]. The deep level emissions labeled E1
Fig. 3. TEM images of SNAB matrix and samples doped with CdTe and treated with post thermal annealing at 555 °C for (a) 0 and (b) 10 h. These images illustrate the morphology, average diameter size and average d-spacing [1 1 1] of the CdTe USQDs.
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Acknowledgements The authors gratefully acknowledge financial support from the agencies CAPES, FAPEMIG and MCT/CNPq. AFM Shimadzu at the Institute of Physics (INFIS) supported by the grant ‘‘PróEquipamentos’’ from the Brazilian Agency CAPES. References
Fig. 4. PL spectra of (a) the SNAB glass matrix and samples doped with CdTe with post thermal annealing at 555 °C for 0, 2 and 10 h. (b) Different recombination processes, where the emissions from CdTe USQDs are identified. The Eexn (2.34 eV), ESDL (2.19 eV), E1 (1.65 eV) and E2 (1.39 eV) emissions refer to the excitonic emission and defect levels present in the CdTe USQDs. The non-radiative processes are also indicated. (c) Ratio of the defect to the excitonic level emission intensities, (ESDL/Eexn), E1/Eexn and E2/Eexn with increasing thermal annealing times. (d) Ratio of the surface defect to the excitonic level emission intensities (ESDL/Eexn) with increasing thermal annealing times.
of CdTe USQDs and the luminescent intensities arising from surface defects and VCd due to the heat treatment time. 4. Conclusions We grew CdTe USQDs in a SNAB (40SiO2 30Na2CO3 1Al2O3 29B2O3 (mol%)) glass matrix by the fusion method and controlled the growth of the CdTe USQDs by varying thermal annealing time. Longer thermal annealing times caused refd shift of the PL and OA bands, indicating growth of USQDs. The average size of the USQDs in the AFM and TEM images confirmed this phenomenon and was in agreement with the size estimated by the EMA model. We also observed that longer thermal annealing time favored decreased luminescence corresponding to the deep level of the VCd within the CdTe USQDs. This was caused in turn by the diffusion of VCd to the surface of CdTe USQDs. Moreover, the decrease of the surface defect level luminescence was observed confirming the increases of size NCs. Therefore, we also controlled VCd diffusion as a function of thermal annealing time. We hope these results stimulate further research into the growth of several types USQDs with desirable properties for optoelectronic device applications.
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