Fig. 3.14. Left: CL from sputter deposited BN:Tb and pure BN. Right: CL spectra of BN and

Maqbool et al. (Maqbool, Richardson et al. 2005) investigated the luminescence of Pr in BN.

**Wavelength (nm)**

**Relative Intensity(a.u)**

**Transition Assignment**

 3 BN transition 415 372 4 3P0 → 3H4 492 973 6 3P1→ 3H5 544 752 8 3P0 → 3H6 628 842 9 3P0 → 3F2 651 748 10 3P0 → 3F3 713 257 11 3P0 → 3F4 736 289

Amorphous Beryllium Nitride, produced by reactive sputtering of Be in nitrogen, was

PL and CL excitation was achieved with 532 nm photons and 10 keV electrons, respectively. Light emission is due to the Sm3+ ions and correspond to the following transitions (see inset on the left in Fig. 3.15): 4G5/2 → 6H5/2 (A), 4G5/2 → 6H7/2 (B), 4G5/2 → 6H9/2 (C), and 4G5/2 → 6H11/2 (D). A photograph illustrating the cathodoluminescence image of the Sm-doped a-

BN:Tb, as deposited and heated to 900 K. At large Tb concentrations, the blue BN

Fig. 3.13. CL from a-GaN: Er.

Maqbool (Maqbool, Richardson et al. 2005) investigated Pr3+ ion luminescence in a-GaN. The data are given in Table 3.7 .

Table 3.7. GaN:Pr

#### **3.2.4 Other amorphous nitrides**

#### **<IP_ADDRESS> BN**

Sputtered boron nitride was shown to be a useful host for RE ions by Maqbool et al. (Maqbool, Richardson et al. 2005) and Allen et al. (Allen S. 2000).

Fig. 3.14. Left: CL from sputter deposited BN:Tb and pure BN. Right: CL spectra of BN and BN:Tb, as deposited and heated to 900 K. At large Tb concentrations, the blue BN luminescence is not visible.

Maqbool et al. (Maqbool, Richardson et al. 2005) investigated the luminescence of Pr in BN. Data are given below in Table 3.8 (Maqbool 2005)

Table 3.8. BN:Pr

#### **<IP_ADDRESS> BeN**

184 Cathodoluminescence

Maqbool (Maqbool, Richardson et al. 2005) investigated Pr3+ ion luminescence in a-GaN.

**Transition Assignment**

3 3P2 → 3H4 418 559

4 3P0 → 3H4 493 5687

6 3P1→ 3H5 532 1895

8 3P0 → 3H6 621 4015

9 3P0 → 3F2 650 4227

10 3P0 → 3F3 713 1508

11 3P0 → 3F4 736 1627

Sputtered boron nitride was shown to be a useful host for RE ions by Maqbool et al.

(Maqbool, Richardson et al. 2005) and Allen et al. (Allen S. 2000).

**Wavelength (nm)**

**Relative Intensity(a.u)**

Fig. 3.13. CL from a-GaN: Er.

The data are given in Table 3.7 .

**Material Transition**

Table 3.7. GaN:Pr

**<IP_ADDRESS> BN**

**3.2.4 Other amorphous nitrides**

**Number**

Amorphous Beryllium Nitride, produced by reactive sputtering of Be in nitrogen, was studied with samarium (Zanatta, Richardson et al. 2007).

PL and CL excitation was achieved with 532 nm photons and 10 keV electrons, respectively. Light emission is due to the Sm3+ ions and correspond to the following transitions (see inset on the left in Fig. 3.15): 4G5/2 → 6H5/2 (A), 4G5/2 → 6H7/2 (B), 4G5/2 → 6H9/2 (C), and 4G5/2 → 6H11/2 (D). A photograph illustrating the cathodoluminescence image of the Sm-doped a-BeN sample is presented in the upper-right corner of Fig. 3.15.

Cathodoluminescence from Amorphous and

**4.1.2 Melting and boiling points**

availability of both d and s electrons.

**4.1.3 Ionization energies**

**4.1.4 Electronic configuration**

transition metals.

**4.1.5 Oxidation states**

orbitals are relatively close for atoms in this row.

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 187

The melting points and the molar enthalpies of fusion of the transition metals are both high in comparison to main group elements. This arises from strong metallic bonding in transition metals which occurs due to delocalization of electrons facilitated by the

In moving across the series of metals from scandium to zinc a small change in the values of the first and second ionization energies is observed. This is due to the build-up of electrons in the immediately underlying d-sub-shells that efficiently shields the 4s electrons from the nucleus and minimizing the increase in effective nuclear charge from element to element. The increases in third and fourth ionization energy values are more rapid. However, the trends in these values show the usual discontinuity half way along the series. The reason is that the five d electrons are all unpaired, in singly occupied orbitals. When the sixth and subsequent electrons enter, the electrons have to share the already occupied orbitals resulting in inter-electron repulsions, which would require less energy to remove an electron. Hence, the third ionization energy curve for the last five elements is identical in

shape to the curve for the first five elements, but displaced upwards by 580 kJ mol-1.

The electronic configuration of the atoms of the first row transition elements are basically the same. It can be seen in Table 4.1 that there is a gradual filling of the 3d orbitals across the series starting from scandium. This filling is, however, not regular, since at chromium and copper the population of 3d orbitals increase by the acquisition of an electron from the 4s shell. This illustrates an important generalization about orbital energies of the first row transition series. At chromium, both the 3d and 4s orbitals are occupied, but neither is completely filled in preference to the other. This suggests that the energies of the 3d and 4s

In the case of copper, the 3d level is full, but only one electron occupies the 4s orbital. This suggests that in copper the 3d orbital energy is lower than the 4s orbital. Thus the 3d orbital energy has passed from higher to lower as we move across the period from potassium to zinc. However, the whole question of preference of an atom to adopt a particular electronic configuration is not determined by orbital energy alone. In chromium it can be shown that the 4s orbital energy is still below the 3d which suggests a configuration [Ar]3d44s2. However due to the effect of electronic repulsion between the outer electrons the actual configuration becomes [Ar]3d54s1 where all the electrons in the outer orbitals are unpaired. Table 1 gives some of the physical properties and free atom electronic configuration of

Oxidation states of transition metals are very important to study the spectroscopy and luminescence from these metal ions. The partially filled d-shell electrons play important role in the oxidation states of transition metals. To fully understand the phenomena of

Fig. 3.15. Room-temperature luminescence spectra of a Sm-doped a-BeN film. From(Zanatta, Richardson et al. 2007)

#### **4. Transition metal cathodoluminescence**

Elements in the periodic table that make a bridge between the first two groups and last six groups of the main elements are defined as transition metals (TM). This bridging role has given these elements the name *transition* metals. These metals consist of many elements, however, we shall limit our discussion to a few of them, mainly 22Ti, 24Cr, 39Y, and 74W . These metals are specially characterized by their common feature: partial occupancy of the d-shells. This partially filled *d* shell configuration assigns special properties to these elements, including their strong suitability for use as optical and luminescent materials. To study their optical and luminescent properties it will be useful to understand some of their basic characteristics. Housecroft (Housecroft et al 2007) and Mackay (Mackay 1996) reported some of the basic characteristics of transition metals. The discussion of these characteristics in section 4.1 follows Housecroft and Mackay.

#### **4.1 Characteristics of transition metals**

#### **4.1.1 Density and metallic radii**

The transition elements are much denser than the s-block elements and show a gradual increase in density from scandium to copper. This trend in density can be explained by the small and irregular decrease in metallic radii coupled with the relative increase in atomic mass.

#### **4.1.2 Melting and boiling points**

186 Cathodoluminescence

Fig. 3.15. Room-temperature luminescence spectra of a Sm-doped a-BeN film. From(Zanatta,

Elements in the periodic table that make a bridge between the first two groups and last six groups of the main elements are defined as transition metals (TM). This bridging role has given these elements the name *transition* metals. These metals consist of many elements, however, we shall limit our discussion to a few of them, mainly 22Ti, 24Cr, 39Y, and 74W . These metals are specially characterized by their common feature: partial occupancy of the d-shells. This partially filled *d* shell configuration assigns special properties to these elements, including their strong suitability for use as optical and luminescent materials. To study their optical and luminescent properties it will be useful to understand some of their basic characteristics. Housecroft (Housecroft et al 2007) and Mackay (Mackay 1996) reported some of the basic characteristics of transition metals. The discussion of these characteristics

The transition elements are much denser than the s-block elements and show a gradual increase in density from scandium to copper. This trend in density can be explained by the small and irregular decrease in metallic radii coupled with the relative increase in atomic mass.

Richardson et al. 2007)

**4. Transition metal cathodoluminescence**

in section 4.1 follows Housecroft and Mackay.

**4.1 Characteristics of transition metals**

**4.1.1 Density and metallic radii**

The melting points and the molar enthalpies of fusion of the transition metals are both high in comparison to main group elements. This arises from strong metallic bonding in transition metals which occurs due to delocalization of electrons facilitated by the availability of both d and s electrons.

#### **4.1.3 Ionization energies**

In moving across the series of metals from scandium to zinc a small change in the values of the first and second ionization energies is observed. This is due to the build-up of electrons in the immediately underlying d-sub-shells that efficiently shields the 4s electrons from the nucleus and minimizing the increase in effective nuclear charge from element to element. The increases in third and fourth ionization energy values are more rapid. However, the trends in these values show the usual discontinuity half way along the series. The reason is that the five d electrons are all unpaired, in singly occupied orbitals. When the sixth and subsequent electrons enter, the electrons have to share the already occupied orbitals resulting in inter-electron repulsions, which would require less energy to remove an electron. Hence, the third ionization energy curve for the last five elements is identical in shape to the curve for the first five elements, but displaced upwards by 580 kJ mol-1.

#### **4.1.4 Electronic configuration**

The electronic configuration of the atoms of the first row transition elements are basically the same. It can be seen in Table 4.1 that there is a gradual filling of the 3d orbitals across the series starting from scandium. This filling is, however, not regular, since at chromium and copper the population of 3d orbitals increase by the acquisition of an electron from the 4s shell. This illustrates an important generalization about orbital energies of the first row transition series. At chromium, both the 3d and 4s orbitals are occupied, but neither is completely filled in preference to the other. This suggests that the energies of the 3d and 4s orbitals are relatively close for atoms in this row.

In the case of copper, the 3d level is full, but only one electron occupies the 4s orbital. This suggests that in copper the 3d orbital energy is lower than the 4s orbital. Thus the 3d orbital energy has passed from higher to lower as we move across the period from potassium to zinc. However, the whole question of preference of an atom to adopt a particular electronic configuration is not determined by orbital energy alone. In chromium it can be shown that the 4s orbital energy is still below the 3d which suggests a configuration [Ar]3d44s2. However due to the effect of electronic repulsion between the outer electrons the actual configuration becomes [Ar]3d54s1 where all the electrons in the outer orbitals are unpaired. Table 1 gives some of the physical properties and free atom electronic configuration of transition metals.

#### **4.1.5 Oxidation states**

Oxidation states of transition metals are very important to study the spectroscopy and luminescence from these metal ions. The partially filled d-shell electrons play important role in the oxidation states of transition metals. To fully understand the phenomena of

Cathodoluminescence from Amorphous and

period since it has five unpaired electrons.

to have 5 d-electrons.

these metal ions.

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 189

Scandium is one of the two elements in the first transition metal period which has only one oxidation state (zinc is the other, with an oxidation state of +2). All the other elements have at least two different oxidation states. Manganese, which is in the middle of the period, has the highest number of oxidation states, and indeed the highest oxidation state in the whole

It was mentioned previously that both copper and chromium do not follow the general formula for transition metal oxidation states. This is because copper has 9 d-electrons, which would produce 4 paired d-electrons and 1 unpaired d-electron. Since copper is just 1 electron short of having a completely full d-orbital, it steals an electron from the s-orbital, allowing it to have 10 d-electrons. Likewise, chromium has 4 d-electrons, only 1 short of having a half-filled d-orbital, so it steals an electron from the s-orbital, allowing chromium

Luminescence and spectroscopic properties of materials play an important role in optical device fabrication, display technologies and the laser industry; transition metals are not exceptions. The optical properties of transition metal ion solids have been studied for many years. Hidalgo (Hidalgo, Mendez et al. 1998) and Muller (Muller, Zhou et al. 2009) studied that partially occupied d-shells of these elements play important role in the luminescence and spectroscopy of these elements. Thurbide (Thurbide and Aue 2002), Gedam (Gedam, Dhoble et al. 2007), Grinberg (Grinberg, Barzowska et al. 2001) and Lapraz (Lapraz, Iacconi et al. 1991) reported that in many host materials, these partially occupied d-shells give rise to several important technological applications and in particular, production of high resistivity materials and photonic devices through transition-metal doping has been widely used. Maqbool (Maqbool, Wilson et al. 2010) and Martin (Martin, Spalding et al. 2001) showed that thermal activation, the oxygen effect and co-dopants are good tools to obtain high efficiency and improved luminescence from

Recent progress toward nitride-based light-emitting diodes and electroluminescent devices (ELDs) has been made using crystalline and amorphous nitride semiconductors doped with a variety of transition metals: Richardson (Richardson, Van Patten et al. 2002), Maqbool (Maqbool, Main et al. 2010; Maqbool, Wilson et al. 2010) and Caldwell (Caldwell, Martin et al. 2001). The amorphous III-nitride semiconductors are equally important as their crystalline counterpart because the amorphous material can be grown at room temperature with little stress due to lattice mismatch. They may also be more suitable for waveguides and cylindrical and spherical laser cavities because of the elimination of grain boundaries at low temperature growth. Cathodoluminescence, spectroscopy and the effects of various

Thin films of a-AlN:Cr were prepared and deposited on Si (100) substrate by the method of plasma magnetron sputtering at low temperature as described earlier. The x-ray diffraction

**4.2 Luminescence from pure TM doped in nitride semiconductors**

factors on the luminescence of a few TMs are given below.

**4.2.1 Chromium doped in amorphous aluminum nitride**

(XRD) analysis confirmed that the deposited films were amorphous.

oxidation states of transition metals, we have to understand how the unpaired d-orbital electrons bond. There are five orbitals in a d subshell manifold. As the number of unpaired valence electrons increases, the d-orbital increases, the highest oxidation state increases. This is because unpaired valence electrons are unstable and eager to bond with other chemical species. This means that the oxidation states would be the highest in the very middle of the transition metal periods due to the presence of the highest number of unpaired valence electrons. To determine the oxidation state, unpaired d-orbital electrons are added to the 2s-orbital electrons since the 3d-orbital is located before the 4s-orbital in the periodic table. For example: Scandium has one unpaired electron in the d-orbital. It is added to the 2 electrons of the s-orbital and therefore the oxidation state is +3. So that would mathematically look like: 1s electron + 1s electron + 1d electron = 3 total electrons = oxidation state of +3.

The formula for determining oxidation states would be (with the exception of copper and chromium):

 = *Highest Oxidation State for a Transition metal Number of Unpaired d electrons Two s orbital electrons*

Table 4.1. Physical properties and free atom electronic configuration of transition metals.

oxidation states of transition metals, we have to understand how the unpaired d-orbital electrons bond. There are five orbitals in a d subshell manifold. As the number of unpaired valence electrons increases, the d-orbital increases, the highest oxidation state increases. This is because unpaired valence electrons are unstable and eager to bond with other chemical species. This means that the oxidation states would be the highest in the very middle of the transition metal periods due to the presence of the highest number of unpaired valence electrons. To determine the oxidation state, unpaired d-orbital electrons are added to the 2s-orbital electrons since the 3d-orbital is located before the 4s-orbital in the periodic table. For example: Scandium has one unpaired electron in the d-orbital. It is added to the 2 electrons of the s-orbital and therefore the oxidation state is +3. So that would mathematically look like: 1s electron + 1s electron + 1d electron = 3 total electrons

The formula for determining oxidation states would be (with the exception of copper and

*Number of Unpaired d electrons Two s orbital electrons*

**Radius (pm)**

Sc 3 2.99 1541 2831 164 [Ar] 3d14s2 631 +3

Ti 4 4.50 1660 3287 147 [Ar]3d24s2 658 +2, +3, +4

V 5 5.96 1890 3380 135 [Ar]3d34s2 650 +2, +3, +4,

Cr 6 7.20 1857 2670 129 [Ar]3d54s1 653 +2, +3, +6

Mn 7 7.20 1244 1962 137 [Ar]3d54s2 717 +2, +3, +4,

Fe 8 7.86 1535 2750 126 [Ar]3d64s2 759 +2, +3

Co 9 8.90 1495 2870 125 [Ar]3d74s2 758 +2, +3

Ni 10 8.90 1455 2730 125 [Ar]3d84s2 737 +2

Cu 11 8.92 1083 2567 128 [Ar]3d104s1 746 +2

Zn 12 7.14 420 907 137 [Ar]3d104s2 906 +2

Table 4.1. Physical properties and free atom electronic configuration of transition metals.

**Free atom configuration** **ionization energy (kJ mol-1)**

**Oxidation state**

+5

+6, +7

*Highest Oxidation State for a Transition metal*

**B.P. (°C)**

**M.P. (°C)**

= oxidation state of +3.

**Element Group Density**

 =

**(g/cm3)**

chromium):

Scandium is one of the two elements in the first transition metal period which has only one oxidation state (zinc is the other, with an oxidation state of +2). All the other elements have at least two different oxidation states. Manganese, which is in the middle of the period, has the highest number of oxidation states, and indeed the highest oxidation state in the whole period since it has five unpaired electrons.

It was mentioned previously that both copper and chromium do not follow the general formula for transition metal oxidation states. This is because copper has 9 d-electrons, which would produce 4 paired d-electrons and 1 unpaired d-electron. Since copper is just 1 electron short of having a completely full d-orbital, it steals an electron from the s-orbital, allowing it to have 10 d-electrons. Likewise, chromium has 4 d-electrons, only 1 short of having a half-filled d-orbital, so it steals an electron from the s-orbital, allowing chromium to have 5 d-electrons.

#### **4.2 Luminescence from pure TM doped in nitride semiconductors**

Luminescence and spectroscopic properties of materials play an important role in optical device fabrication, display technologies and the laser industry; transition metals are not exceptions. The optical properties of transition metal ion solids have been studied for many years. Hidalgo (Hidalgo, Mendez et al. 1998) and Muller (Muller, Zhou et al. 2009) studied that partially occupied d-shells of these elements play important role in the luminescence and spectroscopy of these elements. Thurbide (Thurbide and Aue 2002), Gedam (Gedam, Dhoble et al. 2007), Grinberg (Grinberg, Barzowska et al. 2001) and Lapraz (Lapraz, Iacconi et al. 1991) reported that in many host materials, these partially occupied d-shells give rise to several important technological applications and in particular, production of high resistivity materials and photonic devices through transition-metal doping has been widely used. Maqbool (Maqbool, Wilson et al. 2010) and Martin (Martin, Spalding et al. 2001) showed that thermal activation, the oxygen effect and co-dopants are good tools to obtain high efficiency and improved luminescence from these metal ions.

Recent progress toward nitride-based light-emitting diodes and electroluminescent devices (ELDs) has been made using crystalline and amorphous nitride semiconductors doped with a variety of transition metals: Richardson (Richardson, Van Patten et al. 2002), Maqbool (Maqbool, Main et al. 2010; Maqbool, Wilson et al. 2010) and Caldwell (Caldwell, Martin et al. 2001). The amorphous III-nitride semiconductors are equally important as their crystalline counterpart because the amorphous material can be grown at room temperature with little stress due to lattice mismatch. They may also be more suitable for waveguides and cylindrical and spherical laser cavities because of the elimination of grain boundaries at low temperature growth. Cathodoluminescence, spectroscopy and the effects of various factors on the luminescence of a few TMs are given below.

#### **4.2.1 Chromium doped in amorphous aluminum nitride**

Thin films of a-AlN:Cr were prepared and deposited on Si (100) substrate by the method of plasma magnetron sputtering at low temperature as described earlier. The x-ray diffraction (XRD) analysis confirmed that the deposited films were amorphous.

Cathodoluminescence from Amorphous and

such results in other materials as well.

can also be used in biomedical applications.

shown in Fig. 4.3.

**4.2.2 Tungsten doped in amorphous aluminum nitride**

films.

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 191

A strong emission has occurred at 702 nm indicated by a sharp peak in Fig. 4.2. This peak corresponds to 4T2 → 4A2 transition. It is clear from the figure that thermal activation enhances the luminescence intensity six times. A small peak at 661 nm is also observed. This peak corresponds to Ho+3 ion, indicating the presence of holmium impurity in the

Figure 4.2 reports that the intensity of the emission is not only strong but an enhancement in the luminescence is possible by various means like thermal activation. The strong intensity enabled us to see the red emission in CL apparatus directly with naked eye when the excitation current is reasonably high. Six times increase in the emitted light intensity makes it possible that a-AlN:Cr can serve as a potential candidate for a laser production at 702 nm and other optical devices applications. Moreover the significant increase in the intensities of luminescence from Cr+3 ions by thermal activation can be explained on the basis of luminescence from the triply ionized chromium ions. Luminescence occurs from Cr+3 ions and not from Cr+2 or Cr+1. During the film deposition it is most likely that some of Al+3 of AlN may be replaced by Cr+3 but there are also chances for imperfections and defects giving rise to Cr+2 and Cr+1during film growth. These ions do not contribute to luminescence. The smaller the number of these ions, more will be Cr+3 ions and hence luminescence will be higher. When these films are activated thermally at a higher temperature then most of Cr+2 and Cr+1 impurities ionize and convert to Cr+3 ions giving a path to enhanced luminescence. Moreover when the films are transferred to the furnace and thermally activated after removed from the deposition chamber, they are exposed to air. Thus oxidation of the surface of the film cannot be ignored. Oxygen enhances the luminescence of TM ions giving rise to the enhanced luminescence after thermal activation of the films. Chen (Chen, Chen et al. 2000), Little and Kordesch (Little and Kordesch 2001) and Suyver (Suyver et al. 2005) have reported

The results show that amorphous AlN:Cr is a promising candidate for its use in optical and photonic devices and communication tools. The strong red-IR emission makes this material a potential candidate for making laser cavities, quantum dots and other wave-guided applications. Due to the high penetration ability of near infrared light in human tissues, it

Another important member of the TM family is Tungsten (W). Tugsten has the highest melting point and lowest vapor pressure of all metals. It has a very high tensile strength. Our investigations revealed that along with other physical properties tungsten can also be used for visible light emission applications. The cathodoluminescence of tungsten shows that it gives a very broad emission under cathode ray excitation. The emission spectrum is so broad that it covers the entire visible range of the electromagnetic spectrum from 350 nm up to 700 nm. However the dominant portion of the spectrum comes in blue region with a peak at 491 nm. Another peak in blue is also observed at 429 nm. Because of the huge portion of the spectrum and its peaks lying in blue the films also looked blue in appearance when directly exposed to the electron beam in CL. This broad CL emission from tungsten is

Figure 4.1 shows the XRD analysis of the a-AlN:Cr films deposited on flat Si(100) substrate. Only one peak can be observed in the film at 69.1° which corresponds to Si(100). No other peak is present in the figure, indicating that the films deposited on flat silicon substrates are amorphous. Thermal activation of the films at 1200 K has not changed the structure of the films.

Fig. 4.1. XRD analysis of the a-AlN:Cr films deposited on flat Si (100) substrate.

Cathodoluminescence (CL) of the Cr-doped a-AlN was obtained to study the suitability of this TM for various applications in optical and display technology. Figure 4.2 shows the CL spectrum of the as-deposited and thermally activated a-AlN:Cr in 600 nm to 750 nm range.

Fig. 4.2. CL spectra of room temperature and thermally activated a-AlN:Cr films.

Figure 4.1 shows the XRD analysis of the a-AlN:Cr films deposited on flat Si(100) substrate. Only one peak can be observed in the film at 69.1° which corresponds to Si(100). No other peak is present in the figure, indicating that the films deposited on flat silicon substrates are amorphous. Thermal activation of the films at 1200 K has not changed the structure of the

Fig. 4.1. XRD analysis of the a-AlN:Cr films deposited on flat Si (100) substrate.

Fig. 4.2. CL spectra of room temperature and thermally activated a-AlN:Cr films.

Cathodoluminescence (CL) of the Cr-doped a-AlN was obtained to study the suitability of this TM for various applications in optical and display technology. Figure 4.2 shows the CL spectrum of the as-deposited and thermally activated a-AlN:Cr in 600 nm to 750 nm range.

films.

A strong emission has occurred at 702 nm indicated by a sharp peak in Fig. 4.2. This peak corresponds to 4T2 → 4A2 transition. It is clear from the figure that thermal activation enhances the luminescence intensity six times. A small peak at 661 nm is also observed. This peak corresponds to Ho+3 ion, indicating the presence of holmium impurity in the films.

Figure 4.2 reports that the intensity of the emission is not only strong but an enhancement in the luminescence is possible by various means like thermal activation. The strong intensity enabled us to see the red emission in CL apparatus directly with naked eye when the excitation current is reasonably high. Six times increase in the emitted light intensity makes it possible that a-AlN:Cr can serve as a potential candidate for a laser production at 702 nm and other optical devices applications. Moreover the significant increase in the intensities of luminescence from Cr+3 ions by thermal activation can be explained on the basis of luminescence from the triply ionized chromium ions. Luminescence occurs from Cr+3 ions and not from Cr+2 or Cr+1. During the film deposition it is most likely that some of Al+3 of AlN may be replaced by Cr+3 but there are also chances for imperfections and defects giving rise to Cr+2 and Cr+1during film growth. These ions do not contribute to luminescence. The smaller the number of these ions, more will be Cr+3 ions and hence luminescence will be higher. When these films are activated thermally at a higher temperature then most of Cr+2 and Cr+1 impurities ionize and convert to Cr+3 ions giving a path to enhanced luminescence. Moreover when the films are transferred to the furnace and thermally activated after removed from the deposition chamber, they are exposed to air. Thus oxidation of the surface of the film cannot be ignored. Oxygen enhances the luminescence of TM ions giving rise to the enhanced luminescence after thermal activation of the films. Chen (Chen, Chen et al. 2000), Little and Kordesch (Little and Kordesch 2001) and Suyver (Suyver et al. 2005) have reported such results in other materials as well.

The results show that amorphous AlN:Cr is a promising candidate for its use in optical and photonic devices and communication tools. The strong red-IR emission makes this material a potential candidate for making laser cavities, quantum dots and other wave-guided applications. Due to the high penetration ability of near infrared light in human tissues, it can also be used in biomedical applications.

#### **4.2.2 Tungsten doped in amorphous aluminum nitride**

Another important member of the TM family is Tungsten (W). Tugsten has the highest melting point and lowest vapor pressure of all metals. It has a very high tensile strength. Our investigations revealed that along with other physical properties tungsten can also be used for visible light emission applications. The cathodoluminescence of tungsten shows that it gives a very broad emission under cathode ray excitation. The emission spectrum is so broad that it covers the entire visible range of the electromagnetic spectrum from 350 nm up to 700 nm. However the dominant portion of the spectrum comes in blue region with a peak at 491 nm. Another peak in blue is also observed at 429 nm. Because of the huge portion of the spectrum and its peaks lying in blue the films also looked blue in appearance when directly exposed to the electron beam in CL. This broad CL emission from tungsten is shown in Fig. 4.3.

Cathodoluminescence from Amorphous and

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 193

Fig. 4.4. The effect of Gd concentration on the luminescence of a-AlN:W films.

Tungsten is added to Holmium (Ho) in order to observe any enhancement in the green luminescence from Ho. Figure 4.5 shows the effect of varying W concentration on Ho. A comparison between amorphous AlN:1Ho1W, AlN:1Ho3W and thermally activated AlN:1Ho1W is given in this figure. It is clear from the figure that there is no effect of increasing W concentration on the Ho luminescence. The W concentration is increased 3 times which made the W peak broad but no enhancement in Ho. Further, if we compare this figure with the pure W spectrum then a shift can be seen in the W emission wavelength just like that happened due to Gd. And we see from the figure that in the presence of Ho, tungsten emits a single blue light with a wavelength of 461 nm rather than emissions at 429 nm and 491 nm in pure tungsten. A possible explanation may be the holmium has affected the energy level distribution in tungsten. Further it can also be given in explanation that this shift of wavelength did not occur in rare-earth elements discussed in previous chapters but the transition element W has suffered a lot from this shift. Since rare-earth elements are favorable for their internal f-f transitions. These f-levels lying inside the other shells and hence any external change barely affect these transitions. However it is not true for transition metals and they can be significantly affected by other dopant impurities. Moreover this figure also gives us the effect of thermal annealing on the luminescence of W and Ho. It can be easily deduced from the figure that thermal annealing has activated Ho luminescence more than W. There is also existence of a ruby impurity, which is considerably

**4.2.4 Interaction of co-doped Ho+3 with AlN:W**

enhanced in luminescence by thermal activation.

Fig. 4.3. CL spectrum of amorphous AlN:W

#### **4.2.3 Interaction of co-doped Gd+3 with AlN:W**

Due to its broad emission W+3 is also able to interact with other ions when co-doped with it. For this purpose it was also co-doped with holmium and gadolinium separately to test for

luminescence enhancement in the green emission from holmium ions or luminescence enhancement in the AlN:W by gadolinium. It was observed that no change has occurred in Ho luminescence by the addition of W. However the luminescence from Gd has enhanced W luminescence. Figure 4.4 shows how the addition of Gd+3 affects the light emission from W+3.

Figure 4.4 shows luminescence from amorphous AlN films co-doped with 1 piece of W and 1-3 pieces of Gd. It is clear from the figure that the addition of Gd has a dominant effect on the luminescence from W. Luminescence from AlN:1W2Gd is 3 times the luminescence from AlN:1W1Gd and luminescence from AlN:1W3Gd is almost 6 times the luminescence from AlN:1W1Gd. This fact shows a huge increase in intensities in the W luminescence by Gd.

Figure 4.4 also reveals some other information. It can be observed that the main peak in blue shifts with the relative increase in Gd concentration. With one piece of Gd the main peak of W appears at 506 nm. With two pieces of Gd added to AlN:W the peak gets bigger but shifts to 534 nm. However with the addition of another piece of Gd the peak gets bigger but shifts back to 523 nm. A possible explanation for this shift in the peaks may be the activation of new radiative energy levels in W by the Gd. Further, the intensity of Gd peak is reduced with higher concentration of Gd. Table 4.2 is giving the increase in the W intensity with the concentration of Gd and also the intensity of Gd peak with the concentration of Gd.

Due to its broad emission W+3 is also able to interact with other ions when co-doped with it. For this purpose it was also co-doped with holmium and gadolinium separately to test for luminescence enhancement in the green emission from holmium ions or luminescence enhancement in the AlN:W by gadolinium. It was observed that no change has occurred in Ho luminescence by the addition of W. However the luminescence from Gd has enhanced W luminescence. Figure 4.4 shows how the addition of Gd+3 affects the light emission from W+3. Figure 4.4 shows luminescence from amorphous AlN films co-doped with 1 piece of W and 1-3 pieces of Gd. It is clear from the figure that the addition of Gd has a dominant effect on the luminescence from W. Luminescence from AlN:1W2Gd is 3 times the luminescence from AlN:1W1Gd and luminescence from AlN:1W3Gd is almost 6 times the luminescence from AlN:1W1Gd. This fact shows a huge increase in intensities in the W luminescence by Gd.

Figure 4.4 also reveals some other information. It can be observed that the main peak in blue shifts with the relative increase in Gd concentration. With one piece of Gd the main peak of W appears at 506 nm. With two pieces of Gd added to AlN:W the peak gets bigger but shifts to 534 nm. However with the addition of another piece of Gd the peak gets bigger but shifts back to 523 nm. A possible explanation for this shift in the peaks may be the activation of new radiative energy levels in W by the Gd. Further, the intensity of Gd peak is reduced with higher concentration of Gd. Table 4.2 is giving the increase in the W intensity with the

concentration of Gd and also the intensity of Gd peak with the concentration of Gd.

Fig. 4.3. CL spectrum of amorphous AlN:W

**4.2.3 Interaction of co-doped Gd+3 with AlN:W**

Fig. 4.4. The effect of Gd concentration on the luminescence of a-AlN:W films.

#### **4.2.4 Interaction of co-doped Ho+3 with AlN:W**

Tungsten is added to Holmium (Ho) in order to observe any enhancement in the green luminescence from Ho. Figure 4.5 shows the effect of varying W concentration on Ho. A comparison between amorphous AlN:1Ho1W, AlN:1Ho3W and thermally activated AlN:1Ho1W is given in this figure. It is clear from the figure that there is no effect of increasing W concentration on the Ho luminescence. The W concentration is increased 3 times which made the W peak broad but no enhancement in Ho. Further, if we compare this figure with the pure W spectrum then a shift can be seen in the W emission wavelength just like that happened due to Gd. And we see from the figure that in the presence of Ho, tungsten emits a single blue light with a wavelength of 461 nm rather than emissions at 429 nm and 491 nm in pure tungsten. A possible explanation may be the holmium has affected the energy level distribution in tungsten. Further it can also be given in explanation that this shift of wavelength did not occur in rare-earth elements discussed in previous chapters but the transition element W has suffered a lot from this shift. Since rare-earth elements are favorable for their internal f-f transitions. These f-levels lying inside the other shells and hence any external change barely affect these transitions. However it is not true for transition metals and they can be significantly affected by other dopant impurities. Moreover this figure also gives us the effect of thermal annealing on the luminescence of W and Ho. It can be easily deduced from the figure that thermal annealing has activated Ho luminescence more than W. There is also existence of a ruby impurity, which is considerably enhanced in luminescence by thermal activation.

Cathodoluminescence from Amorphous and

annealing.

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 195

Cathodoluminescence of cold deposited Yttrium doped AlN films were characterized for CL in the same way described earlier for other materials. Figure 4.6 represents the luminescence from AlN:Y films when exposed to cathode rays in CL assembly. This figure confirms emission from amorphous AlN:Y films in UV , blue and in bluish green regions. There is also a peak in IR which may correspond to Yttrium or possibly a ruby impurity. The wavelengths correspond to these emissions are 360 nm, 421 nm, 518 nm and 705 nm respectively. The emission 705 nm could be from AlN:Y or from Cr+3 The bluish green and blue emissions are the dominant in intensity and that is why the films appear blue to naked eye when exposed to electron beam in CL. However we already know from the previous materials investigations that thermal annealing tremendously enhance the luminescence from Cr+3 and hence we can analyze this peak obtaining the CL after performing thermal

Figure 4.7 gives a comparison of CL spectra from amorphous AlN:Y before and after thermal activation. Films were thermally activated at 900 °C for one hour and then characterized for CL. From the figure it is clear that thermal activation has almost no effect on the Yttrium luminescence at 360 nm, 421 nm and 518 nm. However the peak at 705 nm is strongly enhanced by thermal activation. This enhancement in the peak at 705 nm is more than three and half times the luminescence from the same peak in the film, which is not thermally activated. Further, this peak was less than half the intensity of peak at 518 nm before thermal activation but after activation it is about one and half time more intense than the peak at 518 nm. One can guess at this stage that the peak at 705 is most probably due to

chromium oxide (Ruby) rather than emission from yttrium ions itself.

Fig. 4.6. CL spectrum of amorphous AlN:Y

**4.2.5 Cathodoluminescence and thermal activation of AlN doped with Yttrium (Y)**

Fig. 4.5. The effect of thermal activation and W concentration on Ho luminescence.

Table 4.2. The effect of Gd+3 concentration on the W and Gd emission.

Fig. 4.5. The effect of thermal activation and W concentration on Ho luminescence.

1 2207 208 Reference

2 1318 705 339

3 538 1159 557

**Intensity of W luminescence (a.u)**  **Percent increase in luminescence of W**

**luminescence (a.u)**

Table 4.2. The effect of Gd+3 concentration on the W and Gd emission.

**Gd Concentration Intensity of Gd**

#### **4.2.5 Cathodoluminescence and thermal activation of AlN doped with Yttrium (Y)**

Cathodoluminescence of cold deposited Yttrium doped AlN films were characterized for CL in the same way described earlier for other materials. Figure 4.6 represents the luminescence from AlN:Y films when exposed to cathode rays in CL assembly. This figure confirms emission from amorphous AlN:Y films in UV , blue and in bluish green regions. There is also a peak in IR which may correspond to Yttrium or possibly a ruby impurity. The wavelengths correspond to these emissions are 360 nm, 421 nm, 518 nm and 705 nm respectively. The emission 705 nm could be from AlN:Y or from Cr+3 The bluish green and blue emissions are the dominant in intensity and that is why the films appear blue to naked eye when exposed to electron beam in CL. However we already know from the previous materials investigations that thermal annealing tremendously enhance the luminescence from Cr+3 and hence we can analyze this peak obtaining the CL after performing thermal annealing.

Figure 4.7 gives a comparison of CL spectra from amorphous AlN:Y before and after thermal activation. Films were thermally activated at 900 °C for one hour and then characterized for CL. From the figure it is clear that thermal activation has almost no effect on the Yttrium luminescence at 360 nm, 421 nm and 518 nm. However the peak at 705 nm is strongly enhanced by thermal activation. This enhancement in the peak at 705 nm is more than three and half times the luminescence from the same peak in the film, which is not thermally activated. Further, this peak was less than half the intensity of peak at 518 nm before thermal activation but after activation it is about one and half time more intense than the peak at 518 nm. One can guess at this stage that the peak at 705 is most probably due to chromium oxide (Ruby) rather than emission from yttrium ions itself.

Fig. 4.6. CL spectrum of amorphous AlN:Y

Cathodoluminescence from Amorphous and

Fig. 4.8. Cathodoluminescence spectrum of amorphous AlN:Ti.

fiber cavity comes out as a microring emission from the fiber.

**5.1 Amorphous AlN:Tm (Ti) on optical fiber in a cylindrical and ring geometry**

The broad emission from a-AlN:Ti, as discussed in the previous section, makes it hard to use it directly for any particular wavelength application. However it has a great potential for narrowing the peak by a resonance emission to produce a laser out of Ti+3. The resonance emission to produce a laser will definitely need a cavity. Rather than making a traditional cavity, however, we made a cavity by depositing a-AlN:Ti, around optical fiber and obtained a microlaser with emission wavelength of 780.5 nm. The fiber acts as a cavity when the light emitted from the deposited film circulates around the fiber. Only those light waves will enhance each other whose wavelength is an integral multiple of their path around the fiber.

Figure 5.1 represent thin film deposition and light propagation around optical fiber, where 5.1(a) shows longitudinal view of the fiber containing thin film around it. The dark red region around the fiber is the a-AlN:Ti, film deposited uniformly around the fiber. Figure 5.1(b) is a cross sectional view of the fiber with the film, showing light propagation in whispering gallery mode (WGM), The diameter 's' of this section is 12 µm. The dark red region around the fiber is the a-AlN:Ti film deposited on the fiber. The thickness 'd' of this film is 4 micron. This makes the total diameter of the fiber and the film around it to be 20 micron (D = 2d + s = 20 µm). The white lines in the film represent the propagation of light in the film. The pattern of light propagation is restricted to the a-AlN:Ti film only, without touching the fiber itself. Such arrangement is known as whispering gallery. Figure 5.1(c) shows pump laser coupled to fiber and microring emission. Optical fiber with a-AlN:Ti film was held such that it's longitudinal axis is perpendicular to the pump laser beam direction. The pump laser hits the thin films around the fiber directly, exciting a-AlN:Ti for infrared emission, that generates whispering gallery modes in the fiber. The laser produced in the

**5. Applications and special geometries**

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 197

#### **4.2.6 Titanium doped amorphous aluminum nitride**

Titanium is one of the important members of the TM family. It is lightweight, strong, corrosion resistant metal and the ninth most plentiful amongst all the elements in the earth's crust. This metal, when doped in nitride semiconductors, is proved to be a very good source to make laser cavities and optical devices. The cathodoluminescence of Ti+3 doped in amorphous AlN shows a broad emission over a wide range from 650 nm up to 900 nm with a peak around 760 nm. Figure 4.8 shows the CL spectrum of a-AlN:Ti. The broad emission is clear from the spectrum.

Fig. 4.7. Comparison of the CL spectra of thermally activated and as deposited amorphous

Titanium is one of the important members of the TM family. It is lightweight, strong, corrosion resistant metal and the ninth most plentiful amongst all the elements in the earth's crust. This metal, when doped in nitride semiconductors, is proved to be a very good source to make laser cavities and optical devices. The cathodoluminescence of Ti+3 doped in amorphous AlN shows a broad emission over a wide range from 650 nm up to 900 nm with a peak around 760 nm. Figure 4.8 shows the CL spectrum of a-AlN:Ti. The broad emission is

**4.2.6 Titanium doped amorphous aluminum nitride**

AlN:Y films.

clear from the spectrum.

Fig. 4.8. Cathodoluminescence spectrum of amorphous AlN:Ti.

#### **5. Applications and special geometries**

#### **5.1 Amorphous AlN:Tm (Ti) on optical fiber in a cylindrical and ring geometry**

The broad emission from a-AlN:Ti, as discussed in the previous section, makes it hard to use it directly for any particular wavelength application. However it has a great potential for narrowing the peak by a resonance emission to produce a laser out of Ti+3. The resonance emission to produce a laser will definitely need a cavity. Rather than making a traditional cavity, however, we made a cavity by depositing a-AlN:Ti, around optical fiber and obtained a microlaser with emission wavelength of 780.5 nm. The fiber acts as a cavity when the light emitted from the deposited film circulates around the fiber. Only those light waves will enhance each other whose wavelength is an integral multiple of their path around the fiber.

Figure 5.1 represent thin film deposition and light propagation around optical fiber, where 5.1(a) shows longitudinal view of the fiber containing thin film around it. The dark red region around the fiber is the a-AlN:Ti, film deposited uniformly around the fiber. Figure 5.1(b) is a cross sectional view of the fiber with the film, showing light propagation in whispering gallery mode (WGM), The diameter 's' of this section is 12 µm. The dark red region around the fiber is the a-AlN:Ti film deposited on the fiber. The thickness 'd' of this film is 4 micron. This makes the total diameter of the fiber and the film around it to be 20 micron (D = 2d + s = 20 µm). The white lines in the film represent the propagation of light in the film. The pattern of light propagation is restricted to the a-AlN:Ti film only, without touching the fiber itself. Such arrangement is known as whispering gallery. Figure 5.1(c) shows pump laser coupled to fiber and microring emission. Optical fiber with a-AlN:Ti film was held such that it's longitudinal axis is perpendicular to the pump laser beam direction. The pump laser hits the thin films around the fiber directly, exciting a-AlN:Ti for infrared emission, that generates whispering gallery modes in the fiber. The laser produced in the fiber cavity comes out as a microring emission from the fiber.

Cathodoluminescence from Amorphous and

**0**

**500**

**1000**

**1500**

Intensity (a.u)

4.6 nm.

**2000**

**2500**

**3000**

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 199

Figure 5.3 gives increase in the observed near infrared laser in a-AlN:Ti with the increasing power of Nd:YAG pump laser. The curve shows that no laser action can be achieved in a-AlN:Ti if the power P of excitation laser is less than a threshold value P0 . For P > P0 an almost linear increase in the titanium microlaser emission was observed with increasing excitation laser power. The threshold power P0 was found to be about 23.5 mW. For P < P0 we just observed a fluorescence emission from the fiber with an emission width of 35 nm.

**750 760 770 780 790 800**

 22.5 mW 30 mW

Wavelength (nm)

Fig. 5.2. Laser action at 780.5 nm in a-AlN:Ti thin films around optical fiber. The Equally spaced peaks at 776.4 nm and 785.5 nm are the secondary modes of laser. Mode spacing is

Fig. 5.1. Whispering gallery mode microlaser cavity formation around optical fiber.

Figure 5.2 shows the band narrowing and laser action in a-AlN:Ti films deposited around the optical fiber excited by 532 nm Nd:YAG laser. The power of the excitation laser is varied between 7.5 mW and 30 mW. The figure shows two emission spectra. The broad spectrum showing just the fluorescence emission from the deposited films was obtained when 15 mW of power from the Nd:YAG laser is used for excitation. This broad emission spectrum is about 20 nm wide and no lasing action is observed in this spectrum under 15 mW excitation power. The sharp and very intense emission with narrow bandwidth shows lasing action in the film. The laser is produced in the whispering gallery mode at 780.5 nm. We also observed the secondary modes in this spectrum, confirming the resonances in the lasing action. The primary mode lasing and the secondary mode emission peak show gain of ~ 20 at an input power around 30 mW.

Fig. 5.1. Whispering gallery mode microlaser cavity formation around optical fiber.

at an input power around 30 mW.

Figure 5.2 shows the band narrowing and laser action in a-AlN:Ti films deposited around the optical fiber excited by 532 nm Nd:YAG laser. The power of the excitation laser is varied between 7.5 mW and 30 mW. The figure shows two emission spectra. The broad spectrum showing just the fluorescence emission from the deposited films was obtained when 15 mW of power from the Nd:YAG laser is used for excitation. This broad emission spectrum is about 20 nm wide and no lasing action is observed in this spectrum under 15 mW excitation power. The sharp and very intense emission with narrow bandwidth shows lasing action in the film. The laser is produced in the whispering gallery mode at 780.5 nm. We also observed the secondary modes in this spectrum, confirming the resonances in the lasing action. The primary mode lasing and the secondary mode emission peak show gain of ~ 20 Figure 5.3 gives increase in the observed near infrared laser in a-AlN:Ti with the increasing power of Nd:YAG pump laser. The curve shows that no laser action can be achieved in a-AlN:Ti if the power P of excitation laser is less than a threshold value P0 . For P > P0 an almost linear increase in the titanium microlaser emission was observed with increasing excitation laser power. The threshold power P0 was found to be about 23.5 mW. For P < P0 we just observed a fluorescence emission from the fiber with an emission width of 35 nm.

Fig. 5.2. Laser action at 780.5 nm in a-AlN:Ti thin films around optical fiber. The Equally spaced peaks at 776.4 nm and 785.5 nm are the secondary modes of laser. Mode spacing is 4.6 nm.

Cathodoluminescence from Amorphous and

**5.2 Textiles**

good agreement with the theoretical calculations.

was 2.8 kV and 0.15uA/cm2. Field of View 30mm.

(Carlson, Khan et al. 2011; Wang, Carlson et al. 2011).

**5.3 Thermometery**

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 201

Using n = 156, the outer diameter D = 20 micron, and m = 1.95, equation (2) gives ∆λ = 4.9 nm. From figure 5.2 we find that our experimental results give ∆λ = 4.6, which is in quite

The a-AlN:Ti laser, we produced, is very important from biomedical applications point of view. Along with its uses in optics and photonics this laser can also be very beneficial in health sciences applications, particularly for diagnostic and therapeutic purposes. Researchers have reported that near infrared light with a wavelength between 700 nm and 900 nm has minimum absorption and the greatest penetration in body tissues (Noriyuki, Ohdan et al. 1997; Cerussi, Shah et al. 2006). Our laser produced at 780.5 nm is in this range and hence, can be used for diagnosis of deep tissues abnormality and tumors, and laser

The deposition of amorphous AlN onto textiles was examined by Kordesch and Richardson (Kordesch and Richardson 2003). The non-woven textiles (clean room suit material) is made by thermally binding polymer threads. The threads are bound by many pads so that no loose threads are lost. The material is sensitive to heat, so that it is a good test for the successful deposition of a RE – doped Al coated textile. In this case the dopant was Tb.

Right: Cathodoluminescence from the a-AlN:Ti coated textile. The voltage and current density used

A novel use of a-AlN:Ti thin films has been developed by Richardson and co-workers

While the example given in Figure 5.5 is necessarily photoluminescence because the intention of the study is to determine the heating due to a gold particle in a liquid or biological environment, Er based thermometry could be used in CL and in other environments. Several other RE based thermometry systems are possible(Alden, Omrane et al.; Lai, Feng et al.).

Fig. 5.4. Left: a-AlN:Ti film on a thermally bonded textile pad. Field of view is 3mm.

surgery of deep body tissues due to its high penetration ability in the human body.

Fig. 5.3. Increase in the a-AlN:Ti laser intensity with the pump laser power. A threshold power of 24 mW can be observed.

In micro-ring cavities, the thin a-AlN:Ti film deposited around the optical fiber supports whispering gallery modes. Frolov (Frolov, Fujii et al. 1998) (Frolov, Vardeny et al. 1998) and Maqbool and Kordesch (Maqbool, Main et al. 2010) figured out that the angular momentum mode number (n) for a WGM is given by equation (1), where D is the diameter of the optical fiber with thin film deposited on it, λn is the wavelength of the WGM and m is the index of refraction for the a-AlN film.

$$\mathbf{n} = \mathbf{n} \, D \, \mathbf{m} \,/\,\,\lambda\_{\mathbf{n}} \tag{1}$$

Khoshman and Kordesch (Khoshman and Kordesch 2005) worked on amorphous AlN and found that the refractive index in the near infrared region (780.5 nm) is 1.95. This value of the index of refraction of amorphous AlN is obtained using the films deposited in the same deposition system and the same conditions that we used for our work. Thus, m = 1.95 is the most appropriate value to use in our work. Our results satisfy the given equation for the integers n = 157. Because the film around the optical fiber is thick enough to satisfy the WGM condition (d/D ≥ 0.2), we did not observe waveguide modes. The a-AlN:Ti doped optical fiber was placed in vertical direction so that the ring laser formation occur in a horizontal axis. The resonance wavelength λn should also support the mode separation ∆λ equation, given by;

$$
\lambda \Delta \lambda = \lambda\_{\text{n}} \text{ - } \lambda\_{\text{n} \ast 1} = \lambda^2 \text{ / } \text{m D m} \tag{2}
$$

Using n = 156, the outer diameter D = 20 micron, and m = 1.95, equation (2) gives ∆λ = 4.9 nm. From figure 5.2 we find that our experimental results give ∆λ = 4.6, which is in quite good agreement with the theoretical calculations.

The a-AlN:Ti laser, we produced, is very important from biomedical applications point of view. Along with its uses in optics and photonics this laser can also be very beneficial in health sciences applications, particularly for diagnostic and therapeutic purposes. Researchers have reported that near infrared light with a wavelength between 700 nm and 900 nm has minimum absorption and the greatest penetration in body tissues (Noriyuki, Ohdan et al. 1997; Cerussi, Shah et al. 2006). Our laser produced at 780.5 nm is in this range and hence, can be used for diagnosis of deep tissues abnormality and tumors, and laser surgery of deep body tissues due to its high penetration ability in the human body.

#### **5.2 Textiles**

200 Cathodoluminescence

Threshold Power = 24 mW

**5 10 15 20 25 30**

Fig. 5.3. Increase in the a-AlN:Ti laser intensity with the pump laser power. A threshold

In micro-ring cavities, the thin a-AlN:Ti film deposited around the optical fiber supports whispering gallery modes. Frolov (Frolov, Fujii et al. 1998) (Frolov, Vardeny et al. 1998) and Maqbool and Kordesch (Maqbool, Main et al. 2010) figured out that the angular momentum mode number (n) for a WGM is given by equation (1), where D is the diameter of the optical fiber with thin film deposited on it, λn is the wavelength of the WGM and m is the index of

Khoshman and Kordesch (Khoshman and Kordesch 2005) worked on amorphous AlN and found that the refractive index in the near infrared region (780.5 nm) is 1.95. This value of the index of refraction of amorphous AlN is obtained using the films deposited in the same deposition system and the same conditions that we used for our work. Thus, m = 1.95 is the most appropriate value to use in our work. Our results satisfy the given equation for the integers n = 157. Because the film around the optical fiber is thick enough to satisfy the WGM condition (d/D ≥ 0.2), we did not observe waveguide modes. The a-AlN:Ti doped optical fiber was placed in vertical direction so that the ring laser formation occur in a horizontal axis. The resonance wavelength λn should also support the mode separation ∆λ

Pump Laser Power (mW)

n = π D m / λn (1)

∆λ = λn - λn+1 = λ2 / π D m (2)

**0**

power of 24 mW can be observed.

refraction for the a-AlN film.

equation, given by;

**500**

**1000**

**1500**

Intensity (a.u)

**2000**

**2500**

**3000**

The deposition of amorphous AlN onto textiles was examined by Kordesch and Richardson (Kordesch and Richardson 2003). The non-woven textiles (clean room suit material) is made by thermally binding polymer threads. The threads are bound by many pads so that no loose threads are lost. The material is sensitive to heat, so that it is a good test for the successful deposition of a RE – doped Al coated textile. In this case the dopant was Tb.

Right: Cathodoluminescence from the a-AlN:Ti coated textile. The voltage and current density used was 2.8 kV and 0.15uA/cm2. Field of View 30mm.

Fig. 5.4. Left: a-AlN:Ti film on a thermally bonded textile pad. Field of view is 3mm.

#### **5.3 Thermometery**

A novel use of a-AlN:Ti thin films has been developed by Richardson and co-workers (Carlson, Khan et al. 2011; Wang, Carlson et al. 2011).

While the example given in Figure 5.5 is necessarily photoluminescence because the intention of the study is to determine the heating due to a gold particle in a liquid or biological environment, Er based thermometry could be used in CL and in other environments. Several other RE based thermometry systems are possible(Alden, Omrane et al.; Lai, Feng et al.).

Cathodoluminescence from Amorphous and

*Semiconductor Research* 6(13): 1-8.

art. no.-W3.16.

77(4): 478-479.

Nanocrystalline Nitride Thin Films Doped with Rare Earth and Transition Metals 203

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Caldwell, M. L., A. L. Martin, et al. (2001). "Emission properties of an amorphous AlN: Cr3+

Caldwell, M. L., P. G. Van Patten, et al. (2001). "Visible luminescent activation of amorphous

Carlson, M. T., A. Khan, et al. (2011). "Local Temperature Determination of Optically Excited

Cerussi, A., N. Shah, et al. (2006). "In vivo absorption, scattering, and physiologic properties

Chen, H., K. Y. Chen, et al. (2000). "Band gap engineering in amorphous AlxGa1-xN: Experiment and ab initio calculations." *Applied Physics Letters* 77(8): 1117-1119. Chen, H., K. Gurumurugan, et al. (2000). "Visible and infrared emission from GaN: Er thin

Chen, K. Y. and D. A. Drabold (2002). "First principles molecular dynamics study of amorphous AlxGa1-xN alloys." *Journal Of Applied Physics* 91(12): 9743-9751. Dieke, G. H. (1968). *Spectra and Energy Levels of Rare Earth Ion in Crystals*. New York, Wiley. Dierolf, V., et al., eds. (2011). *Rare-Earth Doping of Advanced Materials for Photonic Applications*. Dimitrova, V. I., P. G. Van Patten, et al. (2001). "Photo-, cathodo-, and electroluminescence

Dimitrova, V. I., P. G. Van Patten, et al. (2000). "Visible emission from electroluminescent

Ebdah, M. (2011). Engineering the Optical, Structural, Electrial and Magnetic Properties of

Frolov, S. V., A. Fujii, et al. (1998). "Cylindrical microlasers and light emitting devices from

Frolov, S. V., Z. V. Vardeny, et al. (1998). "Plastic microring lasers on fibers and wires."

Gedam, S. C., S. J. Dhoble, et al. (2007). "Dy3+ and Mn2+ emission in KMgSO4Cl phosphor."

Grinberg, M., J. Barzowska, et al. (2001). "Inhomogeneous broadening of the dominant Cr3+

Gurumurugan, K., H. Chen, et al. (1999). "Visible cathodoluminescence of Er-doped amorphous AlN thin films." *Applied Physics Letters* 74(20): 3008-3010. Hassan, Z. (1998). Growth, Characterizationation and Fabrication of GaN-based Device Structures. *Physics and Astronomy*. Athens, Ohio University. PhD: 1-158. Hidalgo, P., B. Mendez, et al. (1998). "Luminescence properties of transition-metal-doped

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conducting polymers." *Applied Physics Letters* 72(22): 2811-2813.

sites in LiTaO3 system." *Journal Of Luminescence* 94: 85-90.

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GaSb." *Physical Review B* 57(11): 6479-6484.

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of 58 malignant breast tumors determined by broadband diffuse optical

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devices using an amorphous AlN: Er3+ thin-film phosphor." *Applied Physics Letters*

Oxides and Nitrides of In-Ga-Zn Thin Films Using Nanotechnology. *Physics and*

Metal Ions in Sputtered BN Thin Films, Ohio Univerity.

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Fig. 5.5. Schematic of the Er thermometer. The Er:AlGaN film is deposited onto a substrate. The gold particle on the film surface is excited with laser light through the objective of a microscope. The photoluminescence spectrum of the Er peaks at 540 and 565 nm are used to determine the temperature of the film surrounding the substrate. From (Carlson, Khan et al. 2011), with permission.

#### **6. Acknowledgments**

We would like to thank all of the former members of our research teams who have worked on Rare Earth luminescence in the amorphous nitrides. In particular, we thank Professors Henryk J. Lozykowski, A. Ricardo Zanatta, and Hugh H. Richardson. This work was funded by several agencies, including grants from the Ballistic Missile Defence Organization and the Defence Advanced Research Projects Agency administered by the Office of Naval Research: N00014-96-1782 entitled "Growth, Doping and Contacts from Wide Band Gap Semiconductors" and grant N00014-99-1-0975 entitled "Band-Gap Engineering of the Amorphous In-Ga-Al Nitride Semiconductor Alloys for Luminescent Devices from the Ultraviolet to the Infrared."

#### **7. References**

Fig. 5.5. Schematic of the Er thermometer. The Er:AlGaN film is deposited onto a substrate. The gold particle on the film surface is excited with laser light through the objective of a microscope. The photoluminescence spectrum of the Er peaks at 540 and 565 nm are used to determine the temperature of the film surrounding the substrate. From (Carlson, Khan et al.

We would like to thank all of the former members of our research teams who have worked on Rare Earth luminescence in the amorphous nitrides. In particular, we thank Professors Henryk J. Lozykowski, A. Ricardo Zanatta, and Hugh H. Richardson. This work was funded by several agencies, including grants from the Ballistic Missile Defence Organization and the Defence Advanced Research Projects Agency administered by the Office of Naval Research: N00014-96-1782 entitled "Growth, Doping and Contacts from Wide Band Gap Semiconductors" and grant N00014-99-1-0975 entitled "Band-Gap Engineering of the Amorphous In-Ga-Al Nitride Semiconductor Alloys for Luminescent Devices from the

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**8**

*Algeria*

**Peculiarity of the Cathodoluminescence of Αlpha- Alumina Prepared by Calcination of**

**Gibbsite Powder or Generated by Oxidation**

*Laboratoire des Structures, Propriétés et Interactions Inter Atomiques (LASPI2A)*

Most of metallic materials functioning at high temperature need to have oxidation resistance. This resistance can be achieved when the material develops, through oxidation, an oxide film which acts as a diffusion barrier while keeping a good adherence. In this respect, alpha alumina clearly acts as such. The oxides of aluminum have been the subject of many investigations because of their commercial importance and scientific interest. The thermal stability and optical properties of pure nanometer-sized alumina powder have received much attention because of their intrinsic interest and commercial value. Nanometer-sized alumina powders are widely applied today. One of its applications is in fluorescent lamps due to the absorption of ultraviolet light. In fact, it can also emit the light under excitation with a suitable wavelength. It is important to note that there are many works about alpha alumina using X-ray diffraction, but there is a need for a more detailed structural analysis. To achieve this more exhaustive structural characterization we have

used the Rietveld refinement method and cathodoluminescence (CL) measurements.

CL spectroscopy is widely used as a contactless and relatively nondestructive method to provide microcharacterization of the optical and electronic properties of luminescent materials. Nevertheless, it is used comparatively rarely for the investigation of oxide semiconductor structures. The major advantage of CL spectroscopy in the case of such structures is that most of the anticipated products of oxidation are luminescent, and it is easy to get excitation across the bandgap of any dielectric with readily available electron beam voltages. The emission occurs for all the luminescent mechanisms present in the

Pure α-Al2O3 crystal is colorless and shows little absorption in the ultraviolet–visible (UV– Vis) range. But various impurities (Ti, Mn, Cr, and Fe) even a trace level causes apparent

The colors arise from very minor amounts of impurity (<1% of the Al3+ replaced by other cations) because the Al2O3 structure apparently does not tolerate substitutions. However,

absorptions which are attributed to various emission centers (Jheeta et al., 2006).

**1. Introduction**

material.

**of a Metallic FeCrAl Alloy**

*Centre Universitaire de Khenchela 40000*

Djelloul Abdelkader and Boumaza Abdecharif

### **Peculiarity of the Cathodoluminescence of Αlpha- Alumina Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy**

Djelloul Abdelkader and Boumaza Abdecharif

*Laboratoire des Structures, Propriétés et Interactions Inter Atomiques (LASPI2A) Centre Universitaire de Khenchela 40000 Algeria*

#### **1. Introduction**

Most of metallic materials functioning at high temperature need to have oxidation resistance. This resistance can be achieved when the material develops, through oxidation, an oxide film which acts as a diffusion barrier while keeping a good adherence. In this respect, alpha alumina clearly acts as such. The oxides of aluminum have been the subject of many investigations because of their commercial importance and scientific interest. The thermal stability and optical properties of pure nanometer-sized alumina powder have received much attention because of their intrinsic interest and commercial value. Nanometer-sized alumina powders are widely applied today. One of its applications is in fluorescent lamps due to the absorption of ultraviolet light. In fact, it can also emit the light under excitation with a suitable wavelength. It is important to note that there are many works about alpha alumina using X-ray diffraction, but there is a need for a more detailed structural analysis. To achieve this more exhaustive structural characterization we have used the Rietveld refinement method and cathodoluminescence (CL) measurements.

CL spectroscopy is widely used as a contactless and relatively nondestructive method to provide microcharacterization of the optical and electronic properties of luminescent materials. Nevertheless, it is used comparatively rarely for the investigation of oxide semiconductor structures. The major advantage of CL spectroscopy in the case of such structures is that most of the anticipated products of oxidation are luminescent, and it is easy to get excitation across the bandgap of any dielectric with readily available electron beam voltages. The emission occurs for all the luminescent mechanisms present in the material.

Pure α-Al2O3 crystal is colorless and shows little absorption in the ultraviolet–visible (UV– Vis) range. But various impurities (Ti, Mn, Cr, and Fe) even a trace level causes apparent absorptions which are attributed to various emission centers (Jheeta et al., 2006).

The colors arise from very minor amounts of impurity (<1% of the Al3+ replaced by other cations) because the Al2O3 structure apparently does not tolerate substitutions. However,

Peculiarity of the Cathodoluminescence of Αlpha- Alumina

structure refinement.

equal to ± 2 cm-1.

150 MPa was applied for 5 min.

**3. Results and discussions**

growth mechanisms (T>1470 K).

wavelength range 200-1000 nm was investigated.

Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy 211

The present work was performed on ferritic ODS commercial FeCrAlY alloy PM2000 (20 wt% Cr, 5.8 wt% Al, 0.5 wt% Ti, 0.5 wt% Y2O3). This alloy is an alumina-forming alloy. The specimens, with dimensions 25×25×5 mm3, are cut from a rolled plate. In the case of the PM2000, sample was oxidised in air at 1223 K for 72h. The thermogravimetric method (TG) and differential thermal analysis (DTA) data were recorded under a dry air flow with a heating rate of 10 K/min in a SETARAM TGDTA92–16.18 thermal analyser. TG measurements were corrected for temperature–dependent buoyancy by subtracting the data of a measurement carried out on an inert sample. The crystalline structure of the sample was investigated by XRD using a PANalytical X'Pert Pro MRD diffractometer configured as follows: Cu tube operating at 40 kV and 30 mA (*λ(Kα1)* = 0.15406 nm, *λ(Kα2)*  = 0.15444 nm). The scan rate (2*θ*) was 1°/min at a step size of 0.025°. The data were processed to realize the conditions of the software program Fullprof Suite for the

The FTIR technique was used in the absorbance mode in the 200-4000 cm-1 range. For oxides all bands have characteristic frequencies between 200 and 1000 cm-1. For the FTIR measurements, samples were prepared by grinding the oxide films scraped from the substrate (PM2000). After calcination, 10 to 100 µg of the powder was drawn, then grinded with 23 ±2 mg of CsBr in order to obtain, a pellet of 200-250 µm in thickness. After grinding, the powder was placed in a mould (5 mm diameter) and a cold isostatic pressure (CIP) of

The FTIR spectra are obtained using a Perkin-Elmer spectrometer at resolution of 8 cm-1. For each sample, 120 scans were used. The apparatus is equipped with a system allowing the reduction of the optical course in air in order to minimize the perturbations associated with ambient air (water vapour and CO2). The uncertainty on the position of the various peaks is

The emitted light under electron beam excitation in a UHV system was analyzed through a quartz window with a Jobin Yvon CP 2000 spectrograph and a CCD detector. The

**3.1 Differential thermal analysis and thermal gravimetry analysis (DTA-TG) of gibbsite**  Fig. 1 shows the typical TG–DTA curves of the gibbsite. TG and DTA curves are indicated as dotted and solid lines, respectively. Concerning the dehydration-dehydroxylation process of gibbsite, the dehydration appears to occur in two steps (around 598 and 803 K respectively) at higher temperature. The expected theoretical loss due to dehydration is 34.6%, the experimental loss is 34.3% a little lower. This difference 0.3% is a bit larger than the experimental uncertainty 0.1%, the starting gibbsite may be slightly dehydrated. The last step (803 to 1273 K) gives no thermal event but appears as a continuous mass loss (about 2%) which corresponds to the elimination of residual hydroxyls. The formation of α– alumina occurs between 1473 and 1533 K (MacKenzie et al., 1999). Finally, the structural transformations to well-crystallised α−alumina are described by nucleation (T<1470 K) and

these trace substitutions can cause intense colors. Ruby is red because of its Cr3+ content. Yellow sapphire owes its color to Fe3+. Blue sapphire derives its color from Fe2+-Ti4+ and Fe2+-Fe3+ intervalence charge transfer. Green sapphires contain a mixture of the blue and green colors.

Strong well-known 2*E* →4*A*2 lines of Cr3+ (693 nm) with a long decay time characterize their luminescence spectra. Besides that, much weaker narrow lines are present, which are connected with Cr-pairs and more complicated complexes. The Mn4+ ion is isoelectronic with Cr3+, i.e., both of them have the same electronic structure of the open shell (3d3 configuration). Thus, the spectroscopic properties of α-Al2O3:Mn4+ are similar to those of ruby (α-Al2O3:Cr3+). Octahedral Mn4+ (3d3) would be expected to show the R-line fluorescence characteristic of isoelectronic Cr3+ and in approximately the same region. The dominate defects for the visible emission might be different for α-alumina powders formed by heating any of the hydrates of aluminium to a sufficiently high temperature.

In the presence of lattice defects, extra luminescence emissions can be observed in the ultraviolet (UV) region upon highly energetic excitation. The main intrinsic defects in the αalumina crystals are oxygen vacancies in different charge states: a neutral vacancy, a vacancy capturing one electron (a F+-center), and a vacancy capturing two electrons (a Fcenter) (Kislov et al., 2004; Michizono et al., 2007; Yu et al., 2002). The observed UV spectrum in α- alumina can be deconvoluted into two distinct sub-band components: an F+ center band, located at around 3.8 eV, and a less intense F-center band, located at around 3.0 eV (Brewer et al., 1980; Boumaza & Djelloul, 2010; Boumaza et al., 2010). Depending on the defect introduction method one can create also F2-centers, *F*<sup>2</sup> -centers and <sup>2</sup> *<sup>F</sup>*<sup>2</sup> -centers (double oxygen vacancy with four, three and two trapped electrons respectively). α-Al2O3 crystals with defects in the oxygen sublattice are actively studied as promising storage materials (Kortov & Milman, 1996). In this connection, it is interesting to study luminescence properties of the nanostructured aluminium oxide and compare them with analogous properties of crystalline samples.

In this chapter, we present X-ray diffraction (Rietveld analysis) and CL measurements of αalumina powders formed by calcination of gibbsite or generated by oxidation of a metallic FeCrAl alloy. The peculiarity of the cathodoluminescence under comparable conditions of α- alumina is discussed.

#### **2. Materials and experimental methods**

Gibbsite powder*,* Al2(OH)6, from Prolabo (no 20 984.298) was used. The powder is made of platelet aggregates and was composed of 64.5–67% Al2O3 and max.: 0.01% Fe2O3, 0.02% SO4, 0.002% heavy metals (as Pb), and 1.0% non precipitable by NH4OH (as SO4). The sample experienced an ignition loss of 33–34.5% at 1000 °C and had a purity grade of 99.7%. Its average particle size (20 μm) was due to the agglomeration of crystallites. The specific surface area of the original sample was 0.5 m²/g.

The gibbsite platelets was calcined in ambient atmosphere (pO2 = 0.21 atm) at 1573 K. The cycle was as follows: heating up to an isothermal temperature at 5 K/min, maintaining for 24h at the calcination temperature and fast cooling down to room temperature (air quench). The calcination temperature was maintained for 24 hours to obtain a well-crystallised product.

The present work was performed on ferritic ODS commercial FeCrAlY alloy PM2000 (20 wt% Cr, 5.8 wt% Al, 0.5 wt% Ti, 0.5 wt% Y2O3). This alloy is an alumina-forming alloy. The specimens, with dimensions 25×25×5 mm3, are cut from a rolled plate. In the case of the PM2000, sample was oxidised in air at 1223 K for 72h. The thermogravimetric method (TG) and differential thermal analysis (DTA) data were recorded under a dry air flow with a heating rate of 10 K/min in a SETARAM TGDTA92–16.18 thermal analyser. TG measurements were corrected for temperature–dependent buoyancy by subtracting the data of a measurement carried out on an inert sample. The crystalline structure of the sample was investigated by XRD using a PANalytical X'Pert Pro MRD diffractometer configured as follows: Cu tube operating at 40 kV and 30 mA (*λ(Kα1)* = 0.15406 nm, *λ(Kα2)*  = 0.15444 nm). The scan rate (2*θ*) was 1°/min at a step size of 0.025°. The data were processed to realize the conditions of the software program Fullprof Suite for the structure refinement.

The FTIR technique was used in the absorbance mode in the 200-4000 cm-1 range. For oxides all bands have characteristic frequencies between 200 and 1000 cm-1. For the FTIR measurements, samples were prepared by grinding the oxide films scraped from the substrate (PM2000). After calcination, 10 to 100 µg of the powder was drawn, then grinded with 23 ±2 mg of CsBr in order to obtain, a pellet of 200-250 µm in thickness. After grinding, the powder was placed in a mould (5 mm diameter) and a cold isostatic pressure (CIP) of 150 MPa was applied for 5 min.

The FTIR spectra are obtained using a Perkin-Elmer spectrometer at resolution of 8 cm-1. For each sample, 120 scans were used. The apparatus is equipped with a system allowing the reduction of the optical course in air in order to minimize the perturbations associated with ambient air (water vapour and CO2). The uncertainty on the position of the various peaks is equal to ± 2 cm-1.

The emitted light under electron beam excitation in a UHV system was analyzed through a quartz window with a Jobin Yvon CP 2000 spectrograph and a CCD detector. The wavelength range 200-1000 nm was investigated.

#### **3. Results and discussions**

210 Cathodoluminescence

these trace substitutions can cause intense colors. Ruby is red because of its Cr3+ content. Yellow sapphire owes its color to Fe3+. Blue sapphire derives its color from Fe2+-Ti4+ and Fe2+-Fe3+ intervalence charge transfer. Green sapphires contain a mixture of the blue and

Strong well-known 2*E* →4*A*2 lines of Cr3+ (693 nm) with a long decay time characterize their luminescence spectra. Besides that, much weaker narrow lines are present, which are connected with Cr-pairs and more complicated complexes. The Mn4+ ion is isoelectronic with Cr3+, i.e., both of them have the same electronic structure of the open shell (3d3 configuration). Thus, the spectroscopic properties of α-Al2O3:Mn4+ are similar to those of ruby (α-Al2O3:Cr3+). Octahedral Mn4+ (3d3) would be expected to show the R-line fluorescence characteristic of isoelectronic Cr3+ and in approximately the same region. The dominate defects for the visible emission might be different for α-alumina powders formed

In the presence of lattice defects, extra luminescence emissions can be observed in the ultraviolet (UV) region upon highly energetic excitation. The main intrinsic defects in the αalumina crystals are oxygen vacancies in different charge states: a neutral vacancy, a vacancy capturing one electron (a F+-center), and a vacancy capturing two electrons (a Fcenter) (Kislov et al., 2004; Michizono et al., 2007; Yu et al., 2002). The observed UV spectrum in α- alumina can be deconvoluted into two distinct sub-band components: an F+ center band, located at around 3.8 eV, and a less intense F-center band, located at around 3.0 eV (Brewer et al., 1980; Boumaza & Djelloul, 2010; Boumaza et al., 2010). Depending on the

(double oxygen vacancy with four, three and two trapped electrons respectively). α-Al2O3 crystals with defects in the oxygen sublattice are actively studied as promising storage materials (Kortov & Milman, 1996). In this connection, it is interesting to study luminescence properties of the nanostructured aluminium oxide and compare them with analogous

In this chapter, we present X-ray diffraction (Rietveld analysis) and CL measurements of αalumina powders formed by calcination of gibbsite or generated by oxidation of a metallic FeCrAl alloy. The peculiarity of the cathodoluminescence under comparable conditions of

Gibbsite powder*,* Al2(OH)6, from Prolabo (no 20 984.298) was used. The powder is made of platelet aggregates and was composed of 64.5–67% Al2O3 and max.: 0.01% Fe2O3, 0.02% SO4, 0.002% heavy metals (as Pb), and 1.0% non precipitable by NH4OH (as SO4). The sample experienced an ignition loss of 33–34.5% at 1000 °C and had a purity grade of 99.7%. Its average particle size (20 μm) was due to the agglomeration of crystallites. The specific

The gibbsite platelets was calcined in ambient atmosphere (pO2 = 0.21 atm) at 1573 K. The cycle was as follows: heating up to an isothermal temperature at 5 K/min, maintaining for 24h at the calcination temperature and fast cooling down to room temperature (air quench). The calcination temperature was maintained for 24 hours to obtain a well-crystallised product.

by heating any of the hydrates of aluminium to a sufficiently high temperature.

defect introduction method one can create also F2-centers, *F*<sup>2</sup>

properties of crystalline samples.

**2. Materials and experimental methods**

surface area of the original sample was 0.5 m²/g.

α- alumina is discussed.

green colors.

#### **3.1 Differential thermal analysis and thermal gravimetry analysis (DTA-TG) of gibbsite**

Fig. 1 shows the typical TG–DTA curves of the gibbsite. TG and DTA curves are indicated as dotted and solid lines, respectively. Concerning the dehydration-dehydroxylation process of gibbsite, the dehydration appears to occur in two steps (around 598 and 803 K respectively) at higher temperature. The expected theoretical loss due to dehydration is 34.6%, the experimental loss is 34.3% a little lower. This difference 0.3% is a bit larger than the experimental uncertainty 0.1%, the starting gibbsite may be slightly dehydrated. The last step (803 to 1273 K) gives no thermal event but appears as a continuous mass loss (about 2%) which corresponds to the elimination of residual hydroxyls. The formation of α– alumina occurs between 1473 and 1533 K (MacKenzie et al., 1999). Finally, the structural transformations to well-crystallised α−alumina are described by nucleation (T<1470 K) and growth mechanisms (T>1470 K).

Peculiarity of the Cathodoluminescence of Αlpha- Alumina

packing of gibbsite-like layers.

K; (c)- 1073 K; (d)- 1573 K, for 24 h.

no longer present and only AlO6 octahedron remain.

phase.

Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy 213

recently, Kogure (Kogure, 1999) proposed a hexagonal lattice with *a*=0.49 nm and an undefined *c* length indicating that χ-alumina structure can be regarded as random close

For samples prepared at 1173 K, the κ phase is expected (see Fig. 2d). Contrary to the χ phase, the crystal structure of κ-alumina is well known (see for example Ref. (Ollivier et al., 1997) and references therein). κ-alumina is orthorhombic with the space group pna21 and results in ten independent atoms positions (four Al and six O). The experimental XRD pattern at 1173K is specific of a pure κ-alumina (Fig. 2d). Nevertheless, the presence of remnant χ-phase cannot be excluded as all the χ peaks also appear in κ structure. Note that the experimental XRD patterns show well crystallized phases, in contrast with the χ

Fig. 3 give SEM images and XRD patterns of gibbsite powder (a) after calcination at (b)- 773

When heating up at temperature above 1573K for 24h, gibbsite transforms into α-alumina, the stable structure. The XRD pattern of the 1573K sample (Fig. 6) shows that only α-Al2O3 is present when compared with 42–1468 JCPDS file. In this structure, tetrahedral Al3+ ions are

Fig. 2. XRD patterns: (a) Gibbsite (33-0018 JCPDS file), (b) alumina formed from gibbsite calcined for 24h at 773K (expected χ phase) (JCPDS 13-0373 file), (c) alumina formed from gibbsite calcined for 62h at 1073K, (d) alumina formed from gibbsite calcined for 24h at

1173K (expected κ phase) (JCPDS 88-0107 and 01-1305 file).

Fig. 1. TG and DTA curves of gibbsite.

#### **3.2 Gibbsite**

Gibbsite (γ-Al(OH)3) has monoclinic symmetry (*a*=0.8684, *b*=0.5078, *c*=0.9736 nm, β=94.54°) with the space group P21/n, and the unit cell contains eight Al(OH)3 units (Saalfeld & Wedde, 1974). Gibbsite is characterized by the stacking of two-layer units (AA or BB) of hydroxyl sheets with the sequence ABBAABBA... where hydroxyl sheets of the adjacent Gibbsite layers face the *c* direction (Kogure, 1999). In Fig. 2a, the XRD pattern obtained on the as received gibbsite powder shows a good agreement with the reference XRD pattern (33-0018 JCPDS file).

#### **3.2.1 Phase transitions induced by heat treatment of gibbsite**

When heating up fine-grained gibbsite, most OH groups are eliminated, and various forms of alumina are formed with the sequence: gibbsite →χ-→κ-→α-Al2O3 when temperature increases. In order to study these phases, we performed XRD measurements on four samples prepared from gibbsite calcined for 24h at 773, 1073 (62h), 1173 and 1573K respectively.

According to Fig. 2b, at 773K ,the χ phase is expected. In spite of many investigations since the1950's (Brindley & Choe, 1961; Kogure, 1999; Saalfeld, 1960; Stumpf et al., 1950; Yu et al., 2002), the crystal structure of χ-alumina is still uncertain. Stumpth et al. (Stumpf et al., 1950) assumed a cubic (not spinel) unit cell of lattice parameter *a*=0.795 nm (04-0880 JCPDS file). On the other hand, two hexagonal structures have been suggested, either with the parameters *a*=0.556 nm and *c*=1.344 nm (Saalfeld, 1960) or with *a*=0.557 nm and *c*=0.864 nm (Brindley & Choe, 1961) (13-0373 JCPDS file). The two previous hexagonal unit cells may be described respectively as a stacking of 6 and 4 close-packed oxygen layers, of approximately the same thickness (0.224 and 0.216 nm) as the Al–OH layers in gibbsite (0.212 nm). More

Gibbsite (γ-Al(OH)3) has monoclinic symmetry (*a*=0.8684, *b*=0.5078, *c*=0.9736 nm, β=94.54°) with the space group P21/n, and the unit cell contains eight Al(OH)3 units (Saalfeld & Wedde, 1974). Gibbsite is characterized by the stacking of two-layer units (AA or BB) of hydroxyl sheets with the sequence ABBAABBA... where hydroxyl sheets of the adjacent Gibbsite layers face the *c* direction (Kogure, 1999). In Fig. 2a, the XRD pattern obtained on the as received gibbsite powder shows a good agreement with the reference XRD pattern

When heating up fine-grained gibbsite, most OH groups are eliminated, and various forms of alumina are formed with the sequence: gibbsite →χ-→κ-→α-Al2O3 when temperature increases. In order to study these phases, we performed XRD measurements on four samples prepared from gibbsite calcined for 24h at 773, 1073 (62h), 1173 and 1573K

According to Fig. 2b, at 773K ,the χ phase is expected. In spite of many investigations since the1950's (Brindley & Choe, 1961; Kogure, 1999; Saalfeld, 1960; Stumpf et al., 1950; Yu et al., 2002), the crystal structure of χ-alumina is still uncertain. Stumpth et al. (Stumpf et al., 1950) assumed a cubic (not spinel) unit cell of lattice parameter *a*=0.795 nm (04-0880 JCPDS file). On the other hand, two hexagonal structures have been suggested, either with the parameters *a*=0.556 nm and *c*=1.344 nm (Saalfeld, 1960) or with *a*=0.557 nm and *c*=0.864 nm (Brindley & Choe, 1961) (13-0373 JCPDS file). The two previous hexagonal unit cells may be described respectively as a stacking of 6 and 4 close-packed oxygen layers, of approximately the same thickness (0.224 and 0.216 nm) as the Al–OH layers in gibbsite (0.212 nm). More

**3.2.1 Phase transitions induced by heat treatment of gibbsite**

Fig. 1. TG and DTA curves of gibbsite.

**3.2 Gibbsite**

(33-0018 JCPDS file).

respectively.

recently, Kogure (Kogure, 1999) proposed a hexagonal lattice with *a*=0.49 nm and an undefined *c* length indicating that χ-alumina structure can be regarded as random close packing of gibbsite-like layers.

For samples prepared at 1173 K, the κ phase is expected (see Fig. 2d). Contrary to the χ phase, the crystal structure of κ-alumina is well known (see for example Ref. (Ollivier et al., 1997) and references therein). κ-alumina is orthorhombic with the space group pna21 and results in ten independent atoms positions (four Al and six O). The experimental XRD pattern at 1173K is specific of a pure κ-alumina (Fig. 2d). Nevertheless, the presence of remnant χ-phase cannot be excluded as all the χ peaks also appear in κ structure. Note that the experimental XRD patterns show well crystallized phases, in contrast with the χ phase.

Fig. 3 give SEM images and XRD patterns of gibbsite powder (a) after calcination at (b)- 773 K; (c)- 1073 K; (d)- 1573 K, for 24 h.

When heating up at temperature above 1573K for 24h, gibbsite transforms into α-alumina, the stable structure. The XRD pattern of the 1573K sample (Fig. 6) shows that only α-Al2O3 is present when compared with 42–1468 JCPDS file. In this structure, tetrahedral Al3+ ions are no longer present and only AlO6 octahedron remain.

Fig. 2. XRD patterns: (a) Gibbsite (33-0018 JCPDS file), (b) alumina formed from gibbsite calcined for 24h at 773K (expected χ phase) (JCPDS 13-0373 file), (c) alumina formed from gibbsite calcined for 62h at 1073K, (d) alumina formed from gibbsite calcined for 24h at 1173K (expected κ phase) (JCPDS 88-0107 and 01-1305 file).

Peculiarity of the Cathodoluminescence of Αlpha- Alumina

**Mechanical**

Compressive

Maximum Use

Coefficient of Thermal Expansion

**Thermal**

**Electrical**

Fracture Toughness KIC

Thermal

Dielectric

Dielectric

Dissipation

Volume Resistivity

Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy 215

High purity alumina is usable in both oxidizing and reducing atmospheres to 1925°C. Weight loss in vacuum ranges from 10–7 to 10–6 g/cm2.sec over a temperature range of 1700° to 2000°C. It resists attack by all gases except wet fluorine and is resistant to all common reagents except hydrofluoric acid and phosphoric acid. Elevated temperature attack occurs

The composition of the ceramic body can be changed to enhance particular desirable material characteristics. An example would be additions of chrome oxide or manganese oxide to improve hardness and change color. Other additions can be made to improve the ease and consistency of metal films fired to the ceramic for subsequent brazed and soldered assembly.

 **Units of Measure 94% Al2O3 96% Al2O3 99.5% Al2O3**

in the presence of alkali metal vapors particularly at lower purity levels.

Mechanical, thermal and electrical properties of Al2O3 are summarized in table 1.

 Mechanical, thermal and electrical properties of Al2O3

(http://accuratus.com/alumox.html).

MPa•m1/2 3.5 3.5 4

10–6/°C 8.1 8.2 8.4

Ohm•cm >10<sup>14</sup> >10<sup>14</sup> >10<sup>14</sup>

Fig. 3. SEM images: (a)- Gibbsite; (b)- 773 K; (c)- 1073 K; (d)- 1573 K.

#### **3.2.2 General properties of alumina (Al2O3)**

Aluminum oxide, commonly referred to as alumina, possesses strong ionic interatomic bonding giving rise to it's desirable material characteristics. It can exist in several crystalline phases which all revert to the most stable hexagonal alpha phase at elevated temperatures. This is the phase of particular interest for structural applications and the material available from Accuratus.

The exceptional properties of alumina (Al2O3), such as great hardness, high thermal and chemical stability, and high melting temperature, make it a very attractive material. The crystalline α-Al2O3 phase (corundum or sapphire) is the single stable modification of alumina. The crystalline α-Al2O3 has the band gap Eg≈8.5 eV and is widely used in optical devices. Sapphire doped with chrome (ruby) or titanium is applied as an active medium in laser systems. In microelectronics sapphire is used as a substrate for growing silicon and gallium nitride (GaN). Alumina is a highly radiation-resistant material and is used as a sensitive element in detectors when measuring the ionizing radiation parameters.

Fig. 3. SEM images: (a)- Gibbsite; (b)- 773 K; (c)- 1073 K; (d)- 1573 K.

Aluminum oxide, commonly referred to as alumina, possesses strong ionic interatomic bonding giving rise to it's desirable material characteristics. It can exist in several crystalline phases which all revert to the most stable hexagonal alpha phase at elevated temperatures. This is the phase of particular interest for structural applications and the material available

The exceptional properties of alumina (Al2O3), such as great hardness, high thermal and chemical stability, and high melting temperature, make it a very attractive material. The crystalline α-Al2O3 phase (corundum or sapphire) is the single stable modification of alumina. The crystalline α-Al2O3 has the band gap Eg≈8.5 eV and is widely used in optical devices. Sapphire doped with chrome (ruby) or titanium is applied as an active medium in laser systems. In microelectronics sapphire is used as a substrate for growing silicon and gallium nitride (GaN). Alumina is a highly radiation-resistant material and is used as a

sensitive element in detectors when measuring the ionizing radiation parameters.

**3.2.2 General properties of alumina (Al2O3)**

from Accuratus.

High purity alumina is usable in both oxidizing and reducing atmospheres to 1925°C. Weight loss in vacuum ranges from 10–7 to 10–6 g/cm2.sec over a temperature range of 1700° to 2000°C. It resists attack by all gases except wet fluorine and is resistant to all common reagents except hydrofluoric acid and phosphoric acid. Elevated temperature attack occurs in the presence of alkali metal vapors particularly at lower purity levels.

The composition of the ceramic body can be changed to enhance particular desirable material characteristics. An example would be additions of chrome oxide or manganese oxide to improve hardness and change color. Other additions can be made to improve the ease and consistency of metal films fired to the ceramic for subsequent brazed and soldered assembly.

Mechanical, thermal and electrical properties of Al2O3 are summarized in table 1.

Table 1. Mechanical, thermal and electrical properties of Al2O3 (http://accuratus.com/alumox.html).

Peculiarity of the Cathodoluminescence of Αlpha- Alumina

substrate mater as the film is thinner than that formed at 1123 K.

1223 K for 72 h.

Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy 217

Fig. 4 give SEM images and XRD patterns of PM2000 after oxidation at 1023 K for 76 h and

EDX analysis (Fig. 5) of the sample oxidized at 1123 K indicates that the film mainly consists in aluminium and oxygen elements and very small amount of the substrate constituting elements are observed. As the film thickness decreases, like sample oxidized at 1023 K, the signature of the substrate increases due to interactions of the electrons with the underlying

Fig. 4. SEM images of the outer oxidized surface and XRD patterns of α-Al2O3 obtained by

oxidation of PM2000 at 1023 K for 76 h and 1223 K for 72 h.

#### **3.2.3 Oxidation of PM2000**

The oxidation resistance of high-temperature alloys and metallic coatings is dependent on the formation of a protective surface oxide. In an ideal case, the oxide layer should be highly stable, continuous, slow growing, free from cracks or pores, adherent and coherent. α-Al2O3 is an oxide which comes close to satisfying these requirements; the slow growth rate is related to its highly stoichiometric structure and its large band gap which makes electronic conduction difficult.

One of the most crucial factors in the oxidation of alumina-formers is the temperature, which must be high enough to promote the formation of α-Al2O3 in preference to the less protective transition alumina. Another critical factor is the aluminum content which must be sufficiently high to develop and maintain an alumina layer and prevent subsequent breakaway oxidation.

The addition of chromium to Fe-Al alloys promotes the formation and maintenance of a complete layer of α-Al2O3 by acting as a getter and preventing internal oxidation of the aluminum (Wood, 1970).

Iron and chromium are the major impurities present in Al2O3 scales formed on PM2000. Primarily their oxides formed during the transient stage and were incorporated into the α-Al2O3 scale. Fe segregated to some α-Al2O3 grain boundaries, but not Cr. The Al2O3 scale became progressively purer with oxidation time. It is possible that the Fe in the Al2O3 scale increases the scaling rate and, in particular, enhances lateral growth that causes scale convolution.

The growth of α-Al2O3 scales that form on FeCrAl alloys during high temperature oxidation is generally considered to be controlled by oxygen inward diffusion through oxide grain boundaries (Mennicke et al., 1998; Quaddakkers et al., 1991). Aluminum also diffuses out, which can cause growth within the scale (Golightly et al., 1979). The degree of Al outward transport can be significantly reduced by the presence of reactive elements, such as Y, Hf or Zr (Mennicke et al., 1998; Quaddakkers et al., 1991), which segregate to Al2O3 grain boundaries (Przybylski et al., 1987). However, the extent of outward growth seems to differ appreciably among several reactive-element doped Fe based alloys.

The EDS results of the average Fe and Cr concentrations in the scale as a function of scale thickness from different transmission electron microscopy (TEM) specimens are summarized in Table 2.

Standard parameters for α-Al2O3 are: *a* = 0.4758 nm, *c* = 1.2991 nm.

Table 2. Fe and Cr concentrations in α-Al2O3 scales and effect on lattice parameters (Hou et al., 2004)

The oxidation resistance of high-temperature alloys and metallic coatings is dependent on the formation of a protective surface oxide. In an ideal case, the oxide layer should be highly stable, continuous, slow growing, free from cracks or pores, adherent and coherent. α-Al2O3 is an oxide which comes close to satisfying these requirements; the slow growth rate is related to its highly stoichiometric structure and its large band gap which makes electronic conduction difficult.

One of the most crucial factors in the oxidation of alumina-formers is the temperature, which must be high enough to promote the formation of α-Al2O3 in preference to the less protective transition alumina. Another critical factor is the aluminum content which must be sufficiently high to develop and maintain an alumina layer and prevent subsequent

The addition of chromium to Fe-Al alloys promotes the formation and maintenance of a complete layer of α-Al2O3 by acting as a getter and preventing internal oxidation of the

Iron and chromium are the major impurities present in Al2O3 scales formed on PM2000. Primarily their oxides formed during the transient stage and were incorporated into the α-Al2O3 scale. Fe segregated to some α-Al2O3 grain boundaries, but not Cr. The Al2O3 scale became progressively purer with oxidation time. It is possible that the Fe in the Al2O3 scale increases the

The growth of α-Al2O3 scales that form on FeCrAl alloys during high temperature oxidation is generally considered to be controlled by oxygen inward diffusion through oxide grain boundaries (Mennicke et al., 1998; Quaddakkers et al., 1991). Aluminum also diffuses out, which can cause growth within the scale (Golightly et al., 1979). The degree of Al outward transport can be significantly reduced by the presence of reactive elements, such as Y, Hf or Zr (Mennicke et al., 1998; Quaddakkers et al., 1991), which segregate to Al2O3 grain boundaries (Przybylski et al., 1987). However, the extent of outward growth seems to differ

The EDS results of the average Fe and Cr concentrations in the scale as a function of scale thickness from different transmission electron microscopy (TEM) specimens are

1000 °C, 0.5 h 0.39 107±36 4.12±0.95 4.49±1.67 *a* = 0.495±0.004

1000 °C, 1 h 0.9 191±44 0.34±0.28 1.91±0.49 *a* = 0.475±0.004 1000 °C, 26 h 1.77 186±53 0.24±0.28 0.60±0.32 *c* = 1.347 ± 0.008 1200 °C, 2 h 2.94 291±46 0.3±0.64 0.70±0.57 *a* = 0.471±0.004 1200 °C, 120 h 4-5.5 1546±423 0.27±0.20 0.06±0.06 *c* = 1.323 ± 0.008

Table 2. Fe and Cr concentrations in α-Al2O3 scales and effect on lattice parameters (Hou et

**Average [Cr] (at%)**

**Average [Fe] (at%)**

**Lattice parameter (nm)**

*c* = 1.353 ± 0.008

scaling rate and, in particular, enhances lateral growth that causes scale convolution.

appreciably among several reactive-element doped Fe based alloys.

**Oxide grain size (nm)**

Standard parameters for α-Al2O3 are: *a* = 0.4758 nm, *c* = 1.2991 nm.

**3.2.3 Oxidation of PM2000**

breakaway oxidation.

aluminum (Wood, 1970).

summarized in Table 2.

**Scale thickness (µm)**

**Oxidation condition**

al., 2004)

Fig. 4 give SEM images and XRD patterns of PM2000 after oxidation at 1023 K for 76 h and 1223 K for 72 h.

EDX analysis (Fig. 5) of the sample oxidized at 1123 K indicates that the film mainly consists in aluminium and oxygen elements and very small amount of the substrate constituting elements are observed. As the film thickness decreases, like sample oxidized at 1023 K, the signature of the substrate increases due to interactions of the electrons with the underlying substrate mater as the film is thinner than that formed at 1123 K.

Fig. 4. SEM images of the outer oxidized surface and XRD patterns of α-Al2O3 obtained by oxidation of PM2000 at 1023 K for 76 h and 1223 K for 72 h.

Peculiarity of the Cathodoluminescence of Αlpha- Alumina

**3.2.4 Rietveld refinement of the structures**

Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy 219

The Rietveld refinement of the structures was performed using the WinPlotr/FullProf suite package (Rodríguez-Carvajal, 1993). The peak shape was described by a pseudo-Voigt function, and the background level was modeled using a polynomial function. The profiles were refined using the space group and structure models of Al2O3 ( *R C*3 , JCPDS 46-1212)

The XRD data for the 2*θ* regions between 20° and 80° was used for the refinement. The

It was obvious that the agreement between the experimental data and the simulations was excellent since the Rwp (weighted residual error) factor was small (≤ 14.1%). The Rietveld results (cell parameters, atom position, reliability factors and crystallite size (*D*)) are given in Table 3 and 4. The size of α-alumina crystallites (*D*=36 nm) obtained using PM2000 alloy was smaller than that (*D*=43 nm) obtained using gibbsite precursor. Furthermore, the α-Al2O3 from PM2000 has a greater lattice parameter (*a*=0.4763nm, *c*=1.3047nm) than that of JCPDS file 46-1212 (*a*=0.4758 nm, *c*=1.2991 nm, *c*/*a*=2.730) and a greater *c*/*a* ratio (*c*/*a* =2.739) than that of JCPDS file 46-1212. However, the α-Al2O3 from gibbsite has a smaller lattice parameter (*a*=0.4752nm, *c*=1.2980nm) and a similar c/a ratio (*c*/*a* =2.731) than that of JCPDS file 46-1212. Furthermore, the α-Al2O3 from gibbsite has a greater calculated density (3.996 g/cm3) than that of α-Al2O3 from PM2000 (3.961 g/cm3). Phase transformations are frequently accompanied by microstructural changes. This fact could explain the

(nm) atom Wyck. Site *x y z* Biso occupancy

Table 3. The Rietveld refinement results: α-Al2O3 from gibbsite powder, Rp= 16.7%,

Table 4. The Rietveld refinement results: α-Al2O3 from PM2000, Rwp= 14.1%,

(nm) atom Wyck. Site *x y z* Biso occupancy

**Al** 12c 3. 0 0 0.3521 0.2200 0.3333

**O** 18e .2 0.3065 0 0.2500 0.2400 0.5000

**Al** 12c 3. 0 0 0.3417 4.7208 0.9602

**O** 18e .2 0.3392 0 0.2500 2.7191 1.1186

(in this structure cations occupy the 12c sites and oxygen ions the 18e sites).

observed, Rietveld refined and difference patterns are shown in Figs. 6 and 7.

crystallographic parameters differences between a two α-alumina.

Rwp= 6.93%, Calc. density =3.996 g/cm3, *D=*43 nm.

Calc. density =3.961 g/cm3, *D=*36nm.

Lattice parameters

Lattice parameters

*a=* 0.47637 *c=* 1.30472

*a=* 0.47523 *c=* 1.29805

Fig. 5. SEM images of the outer surface and EDX analyses of α-Al2O3 films obtained by oxidation of PM2000 in air at 1023 K for 76 h (top) and 1123 K for 95 h (bottom).

#### **3.2.4 Rietveld refinement of the structures**

218 Cathodoluminescence

Fig. 5. SEM images of the outer surface and EDX analyses of α-Al2O3 films obtained by oxidation of PM2000 in air at 1023 K for 76 h (top) and 1123 K for 95 h (bottom).

The Rietveld refinement of the structures was performed using the WinPlotr/FullProf suite package (Rodríguez-Carvajal, 1993). The peak shape was described by a pseudo-Voigt function, and the background level was modeled using a polynomial function. The profiles were refined using the space group and structure models of Al2O3 ( *R C*3 , JCPDS 46-1212) (in this structure cations occupy the 12c sites and oxygen ions the 18e sites).

The XRD data for the 2*θ* regions between 20° and 80° was used for the refinement. The observed, Rietveld refined and difference patterns are shown in Figs. 6 and 7.

It was obvious that the agreement between the experimental data and the simulations was excellent since the Rwp (weighted residual error) factor was small (≤ 14.1%). The Rietveld results (cell parameters, atom position, reliability factors and crystallite size (*D*)) are given in Table 3 and 4. The size of α-alumina crystallites (*D*=36 nm) obtained using PM2000 alloy was smaller than that (*D*=43 nm) obtained using gibbsite precursor. Furthermore, the α-Al2O3 from PM2000 has a greater lattice parameter (*a*=0.4763nm, *c*=1.3047nm) than that of JCPDS file 46-1212 (*a*=0.4758 nm, *c*=1.2991 nm, *c*/*a*=2.730) and a greater *c*/*a* ratio (*c*/*a* =2.739) than that of JCPDS file 46-1212. However, the α-Al2O3 from gibbsite has a smaller lattice parameter (*a*=0.4752nm, *c*=1.2980nm) and a similar c/a ratio (*c*/*a* =2.731) than that of JCPDS file 46-1212. Furthermore, the α-Al2O3 from gibbsite has a greater calculated density (3.996 g/cm3) than that of α-Al2O3 from PM2000 (3.961 g/cm3). Phase transformations are frequently accompanied by microstructural changes. This fact could explain the crystallographic parameters differences between a two α-alumina.

Table 3. The Rietveld refinement results: α-Al2O3 from gibbsite powder, Rp= 16.7%, Rwp= 6.93%, Calc. density =3.996 g/cm3, *D=*43 nm.

Table 4. The Rietveld refinement results: α-Al2O3 from PM2000, Rwp= 14.1%, Calc. density =3.961 g/cm3, *D=*36nm.

Peculiarity of the Cathodoluminescence of Αlpha- Alumina

**3.2.5 Fourier transform infrared (FTIR) analysis**

molecular presence in air.

**3.3 Cathodoluminescence**

morphology or microstructure.

Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy 221

Fig. 8 compares the FTIR absorbance spectra of α–Al2O3 obtained after calcination of gibbsite and oxidation of PM2000. For the samples calcined at 1573 K, significant spectroscopic bands at ~640, ~594, ~447 cm-1 and ~386 cm-1 appear which are identified to be the characteristic absorption bands of α-Al2O3 (Barker, 1963). This is in good agreement with XRD observations. Common bands exist in all cases, such as the broad OH band centered around 3420 cm−1, and the 1640 cm−1 H2O vibration band (Ma et al., 2008). The very high surface area of these materials results in rapid adsorption of water from the atmosphere because the FTIR samples were kept and grinded in air. Three peaks of very weak intensities at 2850 cm-1, 2920 cm-1 and 2960 cm-1 are observed which are due to C-H stretching vibrations of alkane groups. The absorption in ~2356 cm−1 is due to CO2

Fig. 8. FTIR absorption spectra of α-alumina powder prepared by calcination of gibbsite

CL is the phenomenon of light emission from specimens as a result of interaction with an electron beam. In insulating crystals, the origin of the luminescence arises from impurity atoms (e.g. transition metals or rare earths) in the crystal lattice. Using an electron microscope to produce the electron beam, the spatial distribution of luminescent sites can be observed with submicron spatial resolution, and correlated with features of the specimen

The mechanisms for CL are similar to those for photoluminescence, but the energy input or excitation source is that of an electron beam rather than a visible or ultraviolet light beam. When an energetic (keV range) electron beam propagates within a semiconductor or insulator, the primary electrons lose energy by the creation of electron-hole pairs. These

(black line) or generated by oxidation of PM2000 (red line).

Fig. 6. The observed and calculated diffraction patterns of α-alumina powder prepared by calcination at 1573K for 24h of gibbsite powder. Vertical bars indicate the calculated position of the Bragg peaks. The blue curves in the bottom correspond to the differences between experimental and calculated profiles.

Fig. 7. The observed and calculated diffraction patterns of α-alumina obtained by oxidation of PM2000 at 1223K for 72h.

#### **3.2.5 Fourier transform infrared (FTIR) analysis**

220 Cathodoluminescence

Fig. 6. The observed and calculated diffraction patterns of α-alumina powder prepared by calcination at 1573K for 24h of gibbsite powder. Vertical bars indicate the calculated position of the Bragg peaks. The blue curves in the bottom correspond to the differences between

Fig. 7. The observed and calculated diffraction patterns of α-alumina obtained by oxidation

experimental and calculated profiles.

of PM2000 at 1223K for 72h.

Fig. 8 compares the FTIR absorbance spectra of α–Al2O3 obtained after calcination of gibbsite and oxidation of PM2000. For the samples calcined at 1573 K, significant spectroscopic bands at ~640, ~594, ~447 cm-1 and ~386 cm-1 appear which are identified to be the characteristic absorption bands of α-Al2O3 (Barker, 1963). This is in good agreement with XRD observations. Common bands exist in all cases, such as the broad OH band centered around 3420 cm−1, and the 1640 cm−1 H2O vibration band (Ma et al., 2008). The very high surface area of these materials results in rapid adsorption of water from the atmosphere because the FTIR samples were kept and grinded in air. Three peaks of very weak intensities at 2850 cm-1, 2920 cm-1 and 2960 cm-1 are observed which are due to C-H stretching vibrations of alkane groups. The absorption in ~2356 cm−1 is due to CO2 molecular presence in air.

Fig. 8. FTIR absorption spectra of α-alumina powder prepared by calcination of gibbsite (black line) or generated by oxidation of PM2000 (red line).

#### **3.3 Cathodoluminescence**

CL is the phenomenon of light emission from specimens as a result of interaction with an electron beam. In insulating crystals, the origin of the luminescence arises from impurity atoms (e.g. transition metals or rare earths) in the crystal lattice. Using an electron microscope to produce the electron beam, the spatial distribution of luminescent sites can be observed with submicron spatial resolution, and correlated with features of the specimen morphology or microstructure.

The mechanisms for CL are similar to those for photoluminescence, but the energy input or excitation source is that of an electron beam rather than a visible or ultraviolet light beam. When an energetic (keV range) electron beam propagates within a semiconductor or insulator, the primary electrons lose energy by the creation of electron-hole pairs. These

Peculiarity of the Cathodoluminescence of Αlpha- Alumina

centre absorbs at 6 eV and emits at 3 eV (415 nm).

200 to 600 nm at room temperature.

impurities such as OH.

200 nm.

Prepared by Calcination of Gibbsite Powder or Generated by Oxidation of a Metallic FeCrAl Alloy 223

eV (325 nm) while that of the F centers is observed at a lower photon energy region around 3.0 eV (410 nm). Thus the presently observed luminescence from α-alumina and sapphire which is centered at 330 and 420 nm can be attributed to the F+ and F centers. As for the others, it is noted that the luminescence at 250 and 290 nm is also observed in the case of αalumina. The luminescence intensity at these bands was found to be sensitive to thermal annealing at higher temperatures, and then might be attributed to the effect of some