Fig. 2. Effects of preheating on the compression strength of the preheated pellets with different H/M ratios

Fig. 3. Effects of preheating on FeO content of the preheated pellets with H/M=70:30

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 335

As shown in Fig. 4a, the compression strength of roasted pellets increases quickly with the preheating temperatures varying from 800 to 900C and remains almost unchanged when temperature is between 900 and 950C. While the temperature is further increased, the pellet strength decreases slightly. The main reason is that a compact layer is quickly formed and the residual magnetite in the core of pellet is difficult to be oxidized completely. If pellet is roasted in poor oxidative atmosphere at relatively high temperature, low melting point slag phases (such as 2FeO•SiO2) are formed, and result in deteriorating the compression strength of roasted pellet. From Fig. 4b, the strength of the roasted pellets increased gradually with increasing preheating time, the reason for which is that magnetite particles are gradually oxidized into Fe2O3 crystallites and they are recrystallized more completely

As shown in Fig. 4c, the pellet strength increases steadily with the increase of roasting temperature below 1300C. The possible reason is that the Fe3O4 is gradually oxidized into Fe2O3, which leads to Fe3+ diffusion, rearrangement of Fe2O3 crystal lattices and a compact microstructure formed. When the roasting temperature is over 1300C, however, the pellet strength is reduced with rising temperature, the main reason for which is that: on one hand, it is too late for the residual magnetite in the core of preheated pellet to be oxidized adequately at high temperature, so that low melting point slag phases are formed priorly; and on the other hand, part of Fe2O3 crystal grains decomposes at high temperature reversely, so that the structure of the pellet is destroyed to some degree. As far as the roasting time is concerned, as given in Fig. 4d, the pellet strength is significantly improved

with the roasting time in the range of 10–15 min and remains constant after 15 min.

The results mentioned above show that adding a proper proportion of magnetite concentrate into the hematite pellet is able to improve the compression strength of both

Mineral composition, microstructure and Fe2O3 crystallization of fired pellets at various preheating and roasting temperatures were studied by using Leica DMRXE microscope with

Oxidation of magnetite into hematite is the main reaction during preheating. Oxygen readily diffuses into the interior of the porous pellet and reacts with magnetite particles during preheating. Oxidation always occurs firstly on the surface of particles and cracks, and usually a few small spotty or lamellae hematite are formed (Figs. 5, 6). The oxidation

It is shown that the hematite in the pellets involves two types: the one is original hematite (OH), which is the unreacted hematite from the raw hematite concentrate (particle A), and the other, namely secondary hematite (SH), is formed from the oxidation of magnetite

Obvious differences in colour and shape can be observed between the OH and SH grains. For the OH grains, their colours are bright and white, and their shapes are original and with

with the prolonged preheating time.

preheated and roasted pellets.

an automatic image analyzer.

concentrate (particle B).

**2.4 Induration mechanisms of mixed H/M pellet**

**2.4.1 Crystallization behavior of Fe2O3 during preheating**

process will advance towards the core with increasing temperature.

#### **2.3.2 Roasting characteristics**

The effects of preheating and roasting parameters on the compression strength of the finished pellets with different H/M ratios are presented in Fig. 4.

a-preheating time 10 min, roasting temp. 1275C, roasting time 15 min; b-preheating temp. 900C, roasting temp. 1275C, roasting time 15 min; c-preheating temp. 900C, preheating time 10 min, roasting time 15 min; d-preheating temp. 900C, preheating time 10 min, roasting temp. 1275C

Fig. 4. Effects of preheating and roasting parameters on the compression strength of the finished pellet with different H/M ratios

The effects of preheating and roasting parameters on the compression strength of the

800 850<PHONE_NUMBER>

4 6 8 10 12 14 16

1200 1250 1300 1350

10 12 14 16 18 20

Roasting time(min)

Roasting temperature (℃)

Preheating time(min)

Preheating temperature(℃)

 H/M=70:30 H/M=50:50

Fig. 4. Effects of preheating and roasting parameters on the compression strength of the

finished pellets with different H/M ratios are presented in Fig. 4.

d

c

b

a

a-preheating time 10 min, roasting temp. 1275C, roasting time 15 min; b-preheating temp. 900C, roasting temp. 1275C, roasting time 15 min; c-preheating temp. 900C, preheating time 10 min, roasting time 15 min; d-preheating temp. 900C, preheating time 10 min, roasting temp. 1275C

**2.3.2 Roasting characteristics**

finished pellet with different H/M ratios

4000

4500

5000

4000

4500

5000

Compression strength of roasted pellet (N/P)

As shown in Fig. 4a, the compression strength of roasted pellets increases quickly with the preheating temperatures varying from 800 to 900C and remains almost unchanged when temperature is between 900 and 950C. While the temperature is further increased, the pellet strength decreases slightly. The main reason is that a compact layer is quickly formed and the residual magnetite in the core of pellet is difficult to be oxidized completely. If pellet is roasted in poor oxidative atmosphere at relatively high temperature, low melting point slag phases (such as 2FeO•SiO2) are formed, and result in deteriorating the compression strength of roasted pellet. From Fig. 4b, the strength of the roasted pellets increased gradually with increasing preheating time, the reason for which is that magnetite particles are gradually oxidized into Fe2O3 crystallites and they are recrystallized more completely with the prolonged preheating time.

As shown in Fig. 4c, the pellet strength increases steadily with the increase of roasting temperature below 1300C. The possible reason is that the Fe3O4 is gradually oxidized into Fe2O3, which leads to Fe3+ diffusion, rearrangement of Fe2O3 crystal lattices and a compact microstructure formed. When the roasting temperature is over 1300C, however, the pellet strength is reduced with rising temperature, the main reason for which is that: on one hand, it is too late for the residual magnetite in the core of preheated pellet to be oxidized adequately at high temperature, so that low melting point slag phases are formed priorly; and on the other hand, part of Fe2O3 crystal grains decomposes at high temperature reversely, so that the structure of the pellet is destroyed to some degree. As far as the roasting time is concerned, as given in Fig. 4d, the pellet strength is significantly improved with the roasting time in the range of 10–15 min and remains constant after 15 min.

The results mentioned above show that adding a proper proportion of magnetite concentrate into the hematite pellet is able to improve the compression strength of both preheated and roasted pellets.

#### **2.4 Induration mechanisms of mixed H/M pellet**

Mineral composition, microstructure and Fe2O3 crystallization of fired pellets at various preheating and roasting temperatures were studied by using Leica DMRXE microscope with an automatic image analyzer.

#### **2.4.1 Crystallization behavior of Fe2O3 during preheating**

Oxidation of magnetite into hematite is the main reaction during preheating. Oxygen readily diffuses into the interior of the porous pellet and reacts with magnetite particles during preheating. Oxidation always occurs firstly on the surface of particles and cracks, and usually a few small spotty or lamellae hematite are formed (Figs. 5, 6). The oxidation process will advance towards the core with increasing temperature.

It is shown that the hematite in the pellets involves two types: the one is original hematite (OH), which is the unreacted hematite from the raw hematite concentrate (particle A), and the other, namely secondary hematite (SH), is formed from the oxidation of magnetite concentrate (particle B).

Obvious differences in colour and shape can be observed between the OH and SH grains. For the OH grains, their colours are bright and white, and their shapes are original and with

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 337

By contrast, for the SH grains, they are relatively heavy-coloured, and some residual magnetite domains can be observed. Tabular or massive grains are the main morphology of residual Fe3O4 preheated at 950C (Fig.5), but the grains change into reticular and vein

In comparison with OH grains, the angularities of the SH particles disappear or become unclear during oxidation, and they transform into massive, zonal or columnar-shaped particles. The formation of Fe2O3 microcrystalline junctions between the close SH particles is

The above results suggest that Fe2O3 in SH particles is more active than that in the OH particles. Fe2O3 microcrystalline junctions between SH particles are formed when the pellet is preheated, which is able to improve the strength of the preheated pellet. The OH particles keep their original shapes, and no Fe2O3 microcrystalline junction can be observed at the preheating stage, thus OH has little contribution to the strength of preheated pellets. Therefore, the improvement of oxidative atmosphere during preheating is able to enhance Fe2O3 microcrystalline junctions between the SH particles, the strength of preheated pellet

The crystal morphology of Fe2O3 and microstructure of pellets roasted at different

Fig. 7. OH grains enclosed by or connected with SH grains in the pellet roasted at 1150C

As shown in Fig. 7, a large number of junctions between particles are formed by recrystallization of Fe2O3 at 1150C and the strength of the roasted pellet highly increases. However, the inner structure and shape of the OH particles remain visible, which indicates

shape at 1000C (Fig. 6).

will be improved accordingly.

1-SH particle; 2-OH particle

temperatures are shown in Figs. 7–9.

significantly different from the OH particles at this stage.

**2.4.2 Crystallization behaviour of Fe2O3 during roasting**

discernible angularity. The strip, triangular and rectangular grains can be distinctly observed, and the surface of crystal grain is smooth and the compact inner structure is unchanged. Moreover, the distance between the two close OH particles is large; there is no crystallitic bond formation between them.

1-SH (columnar); 2-OH; 3-residual magnetite (irregular, tabular) Fig. 5. Shapes of SH and OH grains in the pellet preheated at 950C

1-SH (reticular); 2-OH; 3-residual magnetite (vein) Fig. 6. Shapes of SH and OH grains in the pellet preheated at 1000C

discernible angularity. The strip, triangular and rectangular grains can be distinctly observed, and the surface of crystal grain is smooth and the compact inner structure is unchanged. Moreover, the distance between the two close OH particles is large; there is no

crystallitic bond formation between them.

1-SH (columnar); 2-OH; 3-residual magnetite (irregular, tabular)

1-SH (reticular); 2-OH; 3-residual magnetite (vein)

Fig. 6. Shapes of SH and OH grains in the pellet preheated at 1000C

Fig. 5. Shapes of SH and OH grains in the pellet preheated at 950C

By contrast, for the SH grains, they are relatively heavy-coloured, and some residual magnetite domains can be observed. Tabular or massive grains are the main morphology of residual Fe3O4 preheated at 950C (Fig.5), but the grains change into reticular and vein shape at 1000C (Fig. 6).

In comparison with OH grains, the angularities of the SH particles disappear or become unclear during oxidation, and they transform into massive, zonal or columnar-shaped particles. The formation of Fe2O3 microcrystalline junctions between the close SH particles is significantly different from the OH particles at this stage.

The above results suggest that Fe2O3 in SH particles is more active than that in the OH particles. Fe2O3 microcrystalline junctions between SH particles are formed when the pellet is preheated, which is able to improve the strength of the preheated pellet. The OH particles keep their original shapes, and no Fe2O3 microcrystalline junction can be observed at the preheating stage, thus OH has little contribution to the strength of preheated pellets. Therefore, the improvement of oxidative atmosphere during preheating is able to enhance Fe2O3 microcrystalline junctions between the SH particles, the strength of preheated pellet will be improved accordingly.

#### **2.4.2 Crystallization behaviour of Fe2O3 during roasting**

The crystal morphology of Fe2O3 and microstructure of pellets roasted at different temperatures are shown in Figs. 7–9.

1-SH particle; 2-OH particle

Fig. 7. OH grains enclosed by or connected with SH grains in the pellet roasted at 1150C

As shown in Fig. 7, a large number of junctions between particles are formed by recrystallization of Fe2O3 at 1150C and the strength of the roasted pellet highly increases. However, the inner structure and shape of the OH particles remain visible, which indicates

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 339

At 1230C (Fig. 8), Fe2O3 recrystallization is further developed and the interconnection between particles has been enhanced, and individual particles are scarce in pellet. The OH particles have been connected with or even enclosed by SH particles, and the strength of pellet is further increased. However, the inner structure of the OH particles are still compact, with marked difference from the SH particles, which suggests that recrystallization of Fe2O3 in the OH particles is still undeveloped at 1230C, and the junctions between

At 1280C, it can be seen from Fig. 9 that Fe2O3 recrystallization in the OH particles develops satisfactorily; all of hematite particles become porous, the profile of the OH particles almost disappears, and all the particles (both SH and OH) are connected with each other through Fe2O3 recrystallization and form a whole crystal structure; therefore, the strength of pellet is

In summary, the results indicate that Fe2O3 from SH and OH particles have different activities during roasting, which results in the difference of strength formation mechanisms of the mixed H/M concentrates pellets at various roasting temperatures. At lower temperature (1150C), the pellet strength is mainly provided by Fe2O3 recrystallization of the SH particles. However, OH particles can be connected with the SH particles through crystal bridges formed by high activity Fe2O3 of SH particles, or even enclosed by the SH particles.

When the temperature goes up to 1250C, Fe2O3 recrystallization in the SH particles is further developed, and the crystal junctions between the particles become stronger. However, Fe2O3 recrystallization within OH particles does not obviously take place, and few joint is formed between the close OH particles. Pellet strength is mainly provided by Fe2O3

At 1280C, Fe2O3 recrystallization in the SH and OH particles simultaneously occurs, and the joints between the close particles, including OH and SH, are well developed. Particles are connected with each other and the roasted pellet forms a whole crystal structure. The Fe2O3 recrystallization of SH and OH particles plays a crucial role in improving the roasted

The results indicate that, because the activity of Fe2O3 in the SH grains is higher than that in the OH grains, Fe2O3 recrystallization bonds between particles can be enhanced by the SH grains during preheating and roasting, thus the newborn SH is able to improve the pellet strength and decrease the roasting temperature of mixed H/M concentrates pellet. It is the reason that adding a certain proportion of magnetite concentrate is for the enhanced

**3. Fe2O3 recrystallization during the firing of carbon-burdened hematite pellet**

The hematite concentrate is characterized by high total iron grade (67.2% TFe) and less impurities (Table 3). The particle size is 92% undersize 0.074mm and the specific surface

particles are from the Fe2O3 recrystallization of SH particles mainly.

So, OH particles also contribute to the pellet strength to a certain degree.

recrystallization junctions of SH particles.

roasting performance of hematite pellet.

**3.1 Materials and methods**

area is 1629.5cm2/g (Table 4).

further improved.

pellet strength.

**3.1.1 Materials**

that none of Fe2O3 recrystallization takes place in the OH particles at 1150C. The junctions between particles come from Fe2O3 crystallization of the SH particles.

1-SH particle; 2-OH particle

Fig. 8. Recrystallization of Fe2O3 in the pellet roasted at 1230C

Fig. 9. Developed Fe2O3 recrystallization in the pellet roasted at 1280C

At 1230C (Fig. 8), Fe2O3 recrystallization is further developed and the interconnection between particles has been enhanced, and individual particles are scarce in pellet. The OH particles have been connected with or even enclosed by SH particles, and the strength of pellet is further increased. However, the inner structure of the OH particles are still compact, with marked difference from the SH particles, which suggests that recrystallization of Fe2O3 in the OH particles is still undeveloped at 1230C, and the junctions between particles are from the Fe2O3 recrystallization of SH particles mainly.

At 1280C, it can be seen from Fig. 9 that Fe2O3 recrystallization in the OH particles develops satisfactorily; all of hematite particles become porous, the profile of the OH particles almost disappears, and all the particles (both SH and OH) are connected with each other through Fe2O3 recrystallization and form a whole crystal structure; therefore, the strength of pellet is further improved.

In summary, the results indicate that Fe2O3 from SH and OH particles have different activities during roasting, which results in the difference of strength formation mechanisms of the mixed H/M concentrates pellets at various roasting temperatures. At lower temperature (1150C), the pellet strength is mainly provided by Fe2O3 recrystallization of the SH particles. However, OH particles can be connected with the SH particles through crystal bridges formed by high activity Fe2O3 of SH particles, or even enclosed by the SH particles. So, OH particles also contribute to the pellet strength to a certain degree.

When the temperature goes up to 1250C, Fe2O3 recrystallization in the SH particles is further developed, and the crystal junctions between the particles become stronger. However, Fe2O3 recrystallization within OH particles does not obviously take place, and few joint is formed between the close OH particles. Pellet strength is mainly provided by Fe2O3 recrystallization junctions of SH particles.

At 1280C, Fe2O3 recrystallization in the SH and OH particles simultaneously occurs, and the joints between the close particles, including OH and SH, are well developed. Particles are connected with each other and the roasted pellet forms a whole crystal structure. The Fe2O3 recrystallization of SH and OH particles plays a crucial role in improving the roasted pellet strength.

The results indicate that, because the activity of Fe2O3 in the SH grains is higher than that in the OH grains, Fe2O3 recrystallization bonds between particles can be enhanced by the SH grains during preheating and roasting, thus the newborn SH is able to improve the pellet strength and decrease the roasting temperature of mixed H/M concentrates pellet. It is the reason that adding a certain proportion of magnetite concentrate is for the enhanced roasting performance of hematite pellet.

### **3. Fe2O3 recrystallization during the firing of carbon-burdened hematite pellet**

#### **3.1 Materials and methods**

#### **3.1.1 Materials**

338 Recrystallization

that none of Fe2O3 recrystallization takes place in the OH particles at 1150C. The junctions

between particles come from Fe2O3 crystallization of the SH particles.

Fig. 8. Recrystallization of Fe2O3 in the pellet roasted at 1230C

Fig. 9. Developed Fe2O3 recrystallization in the pellet roasted at 1280C

1-SH particle; 2-OH particle

The hematite concentrate is characterized by high total iron grade (67.2% TFe) and less impurities (Table 3). The particle size is 92% undersize 0.074mm and the specific surface area is 1629.5cm2/g (Table 4).

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 341

01234

(Oxygen partial pressure: 20%, airflow: 6L/min, roasting temperature: 1280C, roasting time: 20min)

The effects of roasting temperature on compression strength of pellet are shown in Fig. 11.

Fig. 10. Effects of anthracite dosage on the compression strength of pellet

 With 1% anthracite Without anthracite

**3.2.2 Effects of roasting temperature on compression strength**

(Oxygen partial pressure 20%, airflow: 6L/min, roasting time: 20min)

Fig. 11. Effects of roasting temperature on the compression strength of pellet

Anthracite dosage (%)

1000 1100 1200 1300 1400

Roasting temperature (℃)

1000

0

500

1000

1500

Compression strength of pellet (N/P)

2000

2500

3000

3500

1500

2000

Compression strength of pellet (N/P)

2500

3000

Anthracite powder was used as the material of burdened carbon, of which the specific surface area reaches 6599 cm2/g after grinding.

\*LOI-Loss on ignition

Table 3. Chemical compositions of the materials /%

Table 4. Size distribution and specific surface area of hematite concentrate

#### **3.1.2 Methods**

For each test, 5 kg of hematite concentrate was blended with the given proportion anthracite powder, using 1.25% bentonite as binder. The green balls were prepared in a disc pelletizer with a diameter of 1000 mm, and the green balls with 9~15 mm in diameter were statically dried at 105C in a drying oven for 4 hours.

Firing teats were carried out in an electrically heated shaft furnace. To simulate firing atmosphere, mixed gas of N2/O2 at the given oxygen partial pressure (volume fraction) was pumped into the shaft furnace at a certain flow-rate. The dry balls were charged into heatresistant pot, and then the pot was pushed downwards into the furnace step by step. The pellets were fired at the given temperature for a given period. Afterwards, the roasted pellets were taken out and naturally cooled to ambient temperature. The compression strength of cooled pellets was measured with an LJ-1000 material experimental machine. An average value of 10 pellets is expressed as the compression strength for each test.

#### **3.2 Firing characteristics of carbon-burdened hematite pellet**

#### **3.2.1 Effects of anthracite dosage on compression strength**

The effects of anthracite dosage on the compression strength of the roasted pellets are shown in Fig. 10.

As shown in Fig. 10, the pellet strength with 0.5% anthracite is a little lower than that with no anthracite. When anthracite dosage reaches 1.0%~1.25%, the compression strength goes up to the maximum, however, the strength decreases greatly if the anthracite further increases from 1.5% to 4%. The results indicate appropriate anthracite amount of 1.0%~1.25% may improve the strength of hematite pellet.

Anthracite powder was used as the material of burdened carbon, of which the specific

Materials Fetotal FeO SiO2 CaO MgO Al2O3 LOI\*

Hematite 67.22 0.55 2.17 0.01 0.05 0.55 0.59

Anthracite 1.48 / 6.86 1.51 0.24 5.88 84.08

Bentonite 7.07 / 60.61 0.94 2.2 17.98 10.41

+0.074mm/% 0.074-0.045mm/% -0.045mm/% specific surface area/cm2.g-1

For each test, 5 kg of hematite concentrate was blended with the given proportion anthracite powder, using 1.25% bentonite as binder. The green balls were prepared in a disc pelletizer with a diameter of 1000 mm, and the green balls with 9~15 mm in diameter were statically

Firing teats were carried out in an electrically heated shaft furnace. To simulate firing atmosphere, mixed gas of N2/O2 at the given oxygen partial pressure (volume fraction) was pumped into the shaft furnace at a certain flow-rate. The dry balls were charged into heatresistant pot, and then the pot was pushed downwards into the furnace step by step. The pellets were fired at the given temperature for a given period. Afterwards, the roasted pellets were taken out and naturally cooled to ambient temperature. The compression strength of cooled pellets was measured with an LJ-1000 material experimental machine. An

The effects of anthracite dosage on the compression strength of the roasted pellets are

As shown in Fig. 10, the pellet strength with 0.5% anthracite is a little lower than that with no anthracite. When anthracite dosage reaches 1.0%~1.25%, the compression strength goes up to the maximum, however, the strength decreases greatly if the anthracite further increases from 1.5% to 4%. The results indicate appropriate anthracite amount of

average value of 10 pellets is expressed as the compression strength for each test.

**3.2 Firing characteristics of carbon-burdened hematite pellet 3.2.1 Effects of anthracite dosage on compression strength**

1.0%~1.25% may improve the strength of hematite pellet.

7.79 28.99 63.22 1629.5

Table 4. Size distribution and specific surface area of hematite concentrate

surface area reaches 6599 cm2/g after grinding.

Table 3. Chemical compositions of the materials /%

dried at 105C in a drying oven for 4 hours.

\*LOI-Loss on ignition

**3.1.2 Methods**

shown in Fig. 10.

(Oxygen partial pressure: 20%, airflow: 6L/min, roasting temperature: 1280C, roasting time: 20min) Fig. 10. Effects of anthracite dosage on the compression strength of pellet

#### **3.2.2 Effects of roasting temperature on compression strength**

The effects of roasting temperature on compression strength of pellet are shown in Fig. 11.

(Oxygen partial pressure 20%, airflow: 6L/min, roasting time: 20min) Fig. 11. Effects of roasting temperature on the compression strength of pellet

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 343

0 10 20 30 40 50

Oxygen partial pressure (%)

Fig. 13. Effects of oxygen partial pressure on the compression strength of carbon-burdened

To shed light on the effects of the burdened carbon on the induration of hematite pellet, a test, as shown in Fig. 14, was designed to identify the reduction and decomposition of

As shown in Fig. 14, a cylinder was made by briquetting hematite concentrate firstly, the cylinder bottom is closed, and its inner diameter is 20 mm, the outer diameter is 30 mm. To allow the gas upward injecting into the inner cylinder and penetrating through the anthracite powder layer, many ventages with 0.1 mm diameter were drilled through the

Dry pellets with 2~3 mm in diameter were prepared from hematite concentrate in advance, and then were charged into the surface layer of inner cylinder. The cylinder bottom, anthracite powder layer and pellet layer were separated by inert material of Al2O3 powder

At the beginning of trail, the sample prepared according to Fig. 14 was placed in an electrically heated shaft furnace, and 6L/min N2 with 99.99% purity was pumped into the shaft furnace from the bottom. The sample was taken out and immersed into water immediately after roasted at 1280C for 20 minutes. Subsequently, the FeO content of the

It is shown that FeO content of the cylinder bottom and pellet is 4.35% and 28.69% respectively. FeO content of the pellet is obviously higher than that of the cylinder bottom.

(Airflow: 6L/min, roasting temperature: 1280C, roasting time: 20min, anthracite: 1.0%)

1500

**3.3 Roles of burdened carbon during the roasting**

hematite in carbon-burdened pellet during roasting.

cylinder bottom and pellet was measured respectively.

to avoid their contact with each other.

pellet

cylinder bottom.

2000

2500

Compression strength of pellet (N/P)

3000

As shown in Fig.11, the strength of carbon-burdened pellet is always higher than that of pellet without carbon when roasted at the same temperature; moreover, for carbonburdened hematite pellet, the compression strength increases constantly with the temperature. However, the strength of pellet in the absence of carbon not only doesn't increase markedly until 1250C, but also decreases over 1320C due to the decomposition of Fe2O3. The results imply that the roasting temperature can be decreased and the firing temperature range is enlarged by adding an appropriate amount of anthracite into hematite pellet.

#### **3.2.3 Effects of roasting time on compression strength**

The compression strength of roasted pellet increases gradually with the roasting time prolonging and a maximum strength is obtained at 25 min (in Fig.12), hereafter, the strength almost remains constant.

(Oxygen partial pressure: 20%, airflow: 6L/min, roasting temperature: 1280C, anthracite: 1.0%) Fig. 12. Effects of roasting time on the compression strength of carbon-burdened pellet

#### **3.2.4 Effects of oxygen partial pressure on compression strength**

The compression strength of pellet is susceptible to the change of oxygen partial pressure as shown in Fig.13. The strength of the pellet roasted in N2, that is, in the inert atmosphere, is slightly higher than that in 10% O2. The strength nearly reaches the maximum at 20% O2, and then decreases gradually with increasing the oxygen partial pressure.

(Airflow: 6L/min, roasting temperature: 1280C, roasting time: 20min, anthracite: 1.0%) Fig. 13. Effects of oxygen partial pressure on the compression strength of carbon-burdened

pellet

342 Recrystallization

As shown in Fig.11, the strength of carbon-burdened pellet is always higher than that of pellet without carbon when roasted at the same temperature; moreover, for carbonburdened hematite pellet, the compression strength increases constantly with the temperature. However, the strength of pellet in the absence of carbon not only doesn't increase markedly until 1250C, but also decreases over 1320C due to the decomposition of Fe2O3. The results imply that the roasting temperature can be decreased and the firing temperature range is enlarged by adding an appropriate amount of anthracite into hematite

The compression strength of roasted pellet increases gradually with the roasting time prolonging and a maximum strength is obtained at 25 min (in Fig.12), hereafter, the strength

15 20 25 30

Roasting time (min)

(Oxygen partial pressure: 20%, airflow: 6L/min, roasting temperature: 1280C, anthracite: 1.0%) Fig. 12. Effects of roasting time on the compression strength of carbon-burdened pellet

The compression strength of pellet is susceptible to the change of oxygen partial pressure as shown in Fig.13. The strength of the pellet roasted in N2, that is, in the inert atmosphere, is slightly higher than that in 10% O2. The strength nearly reaches the maximum at 20% O2,

**3.2.4 Effects of oxygen partial pressure on compression strength**

and then decreases gradually with increasing the oxygen partial pressure.

**3.2.3 Effects of roasting time on compression strength**

pellet.

almost remains constant.

3200

2000

2400

Compression strength of pellet (N/P)

2800

#### **3.3 Roles of burdened carbon during the roasting**

To shed light on the effects of the burdened carbon on the induration of hematite pellet, a test, as shown in Fig. 14, was designed to identify the reduction and decomposition of hematite in carbon-burdened pellet during roasting.

As shown in Fig. 14, a cylinder was made by briquetting hematite concentrate firstly, the cylinder bottom is closed, and its inner diameter is 20 mm, the outer diameter is 30 mm. To allow the gas upward injecting into the inner cylinder and penetrating through the anthracite powder layer, many ventages with 0.1 mm diameter were drilled through the cylinder bottom.

Dry pellets with 2~3 mm in diameter were prepared from hematite concentrate in advance, and then were charged into the surface layer of inner cylinder. The cylinder bottom, anthracite powder layer and pellet layer were separated by inert material of Al2O3 powder to avoid their contact with each other.

At the beginning of trail, the sample prepared according to Fig. 14 was placed in an electrically heated shaft furnace, and 6L/min N2 with 99.99% purity was pumped into the shaft furnace from the bottom. The sample was taken out and immersed into water immediately after roasted at 1280C for 20 minutes. Subsequently, the FeO content of the cylinder bottom and pellet was measured respectively.

It is shown that FeO content of the cylinder bottom and pellet is 4.35% and 28.69% respectively. FeO content of the pellet is obviously higher than that of the cylinder bottom.

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 345

FeO content of the roasted pellets with different dosage of anthracite is measured by

As shown in Fig.15, the FeO content of the roasted pellets increases gradually with the increasing the dosage of anthracite. This suggests that the anthracite in pellet benefits the reduction of hematite during roasting. Because of the relatively strong reductive atmosphere under the condition of high dosage of anthracite, a large amount of Fe2O3 is reduced into

The above results that FeO content of the pellet varies with the anthracite dosage are in

01234

(Oxygen partial pressure: 20%, airflow: 6L/min, roasting temperature: 1280C, roasting time: 6min)

FeO content of the pellets roasted at different oxygen contents is measured and the results

As shown in Fig.16, the reduction of hematite mainly occurs at the initial roasting stage. The time, when the maximum FeO content attains to, is shortened with increasing the oxygen partial pressure, which indicates that the increase of oxygen partial pressure has favourable

Fig. 15. Effects of anthracite dosage on FeO content of the carbon-burdened pellet

Anthracite dosage (%)

**3.4.1 Changes of FeO content during the roasting of carbon-burdened pellet**

**3.4 Induration mechanisms of carbon-burdened hematite pellet**

accord with the conclusions obtained from the tests as shown in Fig. 14

chemical analysis and the results are shown in Fig. 15.

**<IP_ADDRESS> Effects of anthracite dosage**

0

are shown in Fig. 16.

**<IP_ADDRESS> Effects of oxygen partial pressure**

effect on oxidization rate of newborn magnetite.

6

12

FeO content of pellet (%)

18

24

Fe3O4.

Fig. 14. Schematic diagram of the test on the role of burdened carbon during roasting

In N2 atmosphere, hematite may decompose into magnetite and release O2 according to formula (1), and FeO content increases accordingly. It is the reaction that the hematite within the cylinder bottom occurs. However, because of being separated by Al2O3 powder, the hematite within the cylinder bottom can't be reduced by anthracite or upwards flowing reductive gases CO/H2, which are produced by gasification of anthracite. Therefore, the increase of FeO content in the cylinder bottom is only caused by the decomposition of hematite.

However, as regards as the hematite within the pellet, on one hand, it may be decomposed into magnetite in N2 gas as same as the hematite within the cylinder bottom; on the other hand, it can be also reduced into magnetite by upwards flowing CO/H2 produced by gasification of anthracite. Therefore, the increase of FeO content of the pellet is caused by both the decomposition and the reduction of hematite, and the latter is more crucial.

The above results identify that the burdened carbon plays the role of reductant during the roasting, and a large number of newborn magnetite are created due to the reductive reaction of hematite by CO/H2, the products of gasification of anthracite.

Of cause, the burdened carbon can release heat via combustion and heat up the inner pellet, which is advantageous for the induration of hematite pellet. The heating function of the burdened carbon resembles the heat release by oxidation of the magnetite concentrate added into hematite pellet.

#### **3.4 Induration mechanisms of carbon-burdened hematite pellet**

#### **3.4.1 Changes of FeO content during the roasting of carbon-burdened pellet**

#### **<IP_ADDRESS> Effects of anthracite dosage**

344 Recrystallization

Fig. 14. Schematic diagram of the test on the role of burdened carbon during roasting

in the cylinder bottom is only caused by the decomposition of hematite.

of hematite by CO/H2, the products of gasification of anthracite.

added into hematite pellet.

In N2 atmosphere, hematite may decompose into magnetite and release O2 according to formula (1), and FeO content increases accordingly. It is the reaction that the hematite within the cylinder bottom occurs. However, because of being separated by Al2O3 powder, the hematite within the cylinder bottom can't be reduced by anthracite or upwards flowing reductive gases CO/H2, which are produced by gasification of anthracite. Therefore, the increase of FeO content

However, as regards as the hematite within the pellet, on one hand, it may be decomposed into magnetite in N2 gas as same as the hematite within the cylinder bottom; on the other hand, it can be also reduced into magnetite by upwards flowing CO/H2 produced by gasification of anthracite. Therefore, the increase of FeO content of the pellet is caused by

The above results identify that the burdened carbon plays the role of reductant during the roasting, and a large number of newborn magnetite are created due to the reductive reaction

Of cause, the burdened carbon can release heat via combustion and heat up the inner pellet, which is advantageous for the induration of hematite pellet. The heating function of the burdened carbon resembles the heat release by oxidation of the magnetite concentrate

both the decomposition and the reduction of hematite, and the latter is more crucial.

FeO content of the roasted pellets with different dosage of anthracite is measured by chemical analysis and the results are shown in Fig. 15.

As shown in Fig.15, the FeO content of the roasted pellets increases gradually with the increasing the dosage of anthracite. This suggests that the anthracite in pellet benefits the reduction of hematite during roasting. Because of the relatively strong reductive atmosphere under the condition of high dosage of anthracite, a large amount of Fe2O3 is reduced into Fe3O4.

The above results that FeO content of the pellet varies with the anthracite dosage are in accord with the conclusions obtained from the tests as shown in Fig. 14

(Oxygen partial pressure: 20%, airflow: 6L/min, roasting temperature: 1280C, roasting time: 6min) Fig. 15. Effects of anthracite dosage on FeO content of the carbon-burdened pellet

#### **<IP_ADDRESS> Effects of oxygen partial pressure**

FeO content of the pellets roasted at different oxygen contents is measured and the results are shown in Fig. 16.

As shown in Fig.16, the reduction of hematite mainly occurs at the initial roasting stage. The time, when the maximum FeO content attains to, is shortened with increasing the oxygen partial pressure, which indicates that the increase of oxygen partial pressure has favourable effect on oxidization rate of newborn magnetite.

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 347

To shed light on the Fe2O3 recrystallization behaviour, the microstructure of carbonburdened hematite pellets at different roasting stage is investigated and the results are

**3.4.2 Fe2O3 recrystallization behaviours of carbon-burdened hematite pellet**

presented in Fig. 17.

(Anthracite: 1.0%, airflow: 6L/min, roasting temperature: 1280C)

Fig. 16. Change of FeO content of the pellets roasted at different oxygen partial pressures for different period

The maximum FeO content decreases when the carbon-burdened hematite pellet is roasted at high oxygen partial pressure. The result shows that the lower the oxygen partial pressure during roasting, the stronger the reductive atmosphere within the pellet; and the reduction of Fe2O3 can be enhanced. At high oxygen partial pressure, the burdened carbon in pellet combusts completely, and more carbon comes into being not CO but CO2, thus reductive atmosphere is weakened, which prevents the hematite from reducing into magnetite.

In oxidative atmosphere, FeO content of the pellet increases firstly and then decreases during roasting, the reason for which is that, at the initial roasting stage, the reduction rate of hematite into magnetite is higher than the oxidation rate of newborn magnetite into hematite, however, accompanied with the consumption of burdened carbon, the reduction rate decreases while the oxidation rate increases, the maximum FeO content attains when the reduction rate equals to the oxidation rate. Subsequently, FeO content decreases alone with the oxidation of magnetite untill the oxidation is complete.

It indicates that some original hematite (OH) can be reduced to magnetite by the burdened carbon firstly, and then the newborn magnetite is oxidized into secondary hematite (SH) again during the roasting. The transformation of OH to magnetite and newborn magnetite to SH changes the route of Fe2O3 recrystallization during the induration of carbon-burdened hematite pellet.

O2

 content 0%(N2) 5% 10% 20% 30% 40% 50%

0 5 10 15 20 25

Fig. 16. Change of FeO content of the pellets roasted at different oxygen partial pressures for

The maximum FeO content decreases when the carbon-burdened hematite pellet is roasted at high oxygen partial pressure. The result shows that the lower the oxygen partial pressure during roasting, the stronger the reductive atmosphere within the pellet; and the reduction of Fe2O3 can be enhanced. At high oxygen partial pressure, the burdened carbon in pellet combusts completely, and more carbon comes into being not CO but CO2, thus reductive atmosphere is weakened, which prevents the hematite from

In oxidative atmosphere, FeO content of the pellet increases firstly and then decreases during roasting, the reason for which is that, at the initial roasting stage, the reduction rate of hematite into magnetite is higher than the oxidation rate of newborn magnetite into hematite, however, accompanied with the consumption of burdened carbon, the reduction rate decreases while the oxidation rate increases, the maximum FeO content attains when the reduction rate equals to the oxidation rate. Subsequently, FeO content decreases alone

It indicates that some original hematite (OH) can be reduced to magnetite by the burdened carbon firstly, and then the newborn magnetite is oxidized into secondary hematite (SH) again during the roasting. The transformation of OH to magnetite and newborn magnetite to SH changes the route of Fe2O3 recrystallization during the induration of carbon-burdened

(Anthracite: 1.0%, airflow: 6L/min, roasting temperature: 1280C)

with the oxidation of magnetite untill the oxidation is complete.

Roasting time (min)

0

reducing into magnetite.

different period

hematite pellet.

5

10

FeO content of pellet (%)

15

20

25

#### **3.4.2 Fe2O3 recrystallization behaviours of carbon-burdened hematite pellet**

To shed light on the Fe2O3 recrystallization behaviour, the microstructure of carbonburdened hematite pellets at different roasting stage is investigated and the results are presented in Fig. 17.

Recrystallization of Fe2O3 During the Induration of Iron Ore Oxidation Pellets 349

When the volatile matter in anthracite is pyrolyzed and gasified during initial roasting stage, the reduction atmosphere is gradually enhanced within the pellet, and a few magnetite grains, shown in Fig. 17a, are formed by the reduction of OH. An apparent colour difference can be observed between the hematite grains and magnetite grains. The hematite presents bright and white colour, while the magnetite is grey. However, the OH particles keep their original shape and discernible angularity. The inner crystal structure of OH is

It can be seen from Fig. 17b, a large amount of newly created magnetite is observed as the roasting process progresses; however, the OH particles keep their original shapes, and there

As shown in Fig. 17c, the content of newborn magnetite decreases, and a few SH, which comes from the oxidation of the newborn magnetite, are formed. In comparison with OH grains, the angularities of the SH particles become unclear, and a few crystal junctions can also be observed between them. In this stage, with the consumption of anthracite, reduction

At the end of roasting, the newborn magnetite has been completely oxidized into SH, and a large number of crystal junctions between particles are formed (as shown in Fig. 17d). Pellet strength is mainly provided by Fe2O3 recrystallization junctions between SH particles and

It can be concluded from the results mentioned above that partial OH particles can be reduced firstly and turn into magnetite particles by the anthracite powder dispersed in the carbon-burdened hematite pellet, however, the newborn magnetite can be subsequently oxidized into SH particles with higher activity, and the route of Fe2O3 recrystallization is changed from the recrystallization of OH particles to the recrystallization bonds among SH particles and OH particles. Therefore, the formation of SH during the roasting of carbonburdened hematite pellet is able to improve the pellet strength and decrease the roasting temperature. It is the reason why adding a certain proportion anthracite is also an effective

Both magnetite and burdened carbon are found to be the favourable techniques for enhancing the induration of hematite pellet. The induration mechanisms of hematite pellet with addition of magnetite concentrate and anthracite powder are revealed by

The crystallization behaviours of Fe2O3 during preheating and roasting of pellets made from mixed hematite/magnetite (H/M) concentrates have been revealed. The results indicate that the strength of pellet is mainly provided by the crystalline connections between the particles during preheating. This occurs because the activity of Fe2O3 from secondary hematite (oxidized from magnetite concentrate) is higher than that from original hematite (from the raw hematite concentrate). In the roasting process, when temperature is lower than 1250C, the strength is mainly provided by the development, connection and growth of Fe2O3 crystalline grains from secondary hematite. Only if the temperature exceeds 1280C does

characterization of Fe2O3 recrystallization rules during the oxidization roasting.

Fe2O3 recrystallization in original hematite grains develop very well.

compact and there is no crystal bond formed between OH particles.

atmosphere is weakened and oxidation rate of SH is accelerated.

way to enhance the roasting performance of hematite pellet.

is no crystal junction observed still.

OH particles.

**4. Conclusions**

Roasting time: a-4min; b-6min; c-8min; d-20min 1-Magnetite; 2-OH;3- SH (Anthracite: 1.0%, Oxygen partial pressure: 20%, airflow: 6L/min)

Fig. 17. Microstructure of the carbon-burdened pellets roasted at 1280C

When the volatile matter in anthracite is pyrolyzed and gasified during initial roasting stage, the reduction atmosphere is gradually enhanced within the pellet, and a few magnetite grains, shown in Fig. 17a, are formed by the reduction of OH. An apparent colour difference can be observed between the hematite grains and magnetite grains. The hematite presents bright and white colour, while the magnetite is grey. However, the OH particles keep their original shape and discernible angularity. The inner crystal structure of OH is compact and there is no crystal bond formed between OH particles.

It can be seen from Fig. 17b, a large amount of newly created magnetite is observed as the roasting process progresses; however, the OH particles keep their original shapes, and there is no crystal junction observed still.

As shown in Fig. 17c, the content of newborn magnetite decreases, and a few SH, which comes from the oxidation of the newborn magnetite, are formed. In comparison with OH grains, the angularities of the SH particles become unclear, and a few crystal junctions can also be observed between them. In this stage, with the consumption of anthracite, reduction atmosphere is weakened and oxidation rate of SH is accelerated.

At the end of roasting, the newborn magnetite has been completely oxidized into SH, and a large number of crystal junctions between particles are formed (as shown in Fig. 17d). Pellet strength is mainly provided by Fe2O3 recrystallization junctions between SH particles and OH particles.

It can be concluded from the results mentioned above that partial OH particles can be reduced firstly and turn into magnetite particles by the anthracite powder dispersed in the carbon-burdened hematite pellet, however, the newborn magnetite can be subsequently oxidized into SH particles with higher activity, and the route of Fe2O3 recrystallization is changed from the recrystallization of OH particles to the recrystallization bonds among SH particles and OH particles. Therefore, the formation of SH during the roasting of carbonburdened hematite pellet is able to improve the pellet strength and decrease the roasting temperature. It is the reason why adding a certain proportion anthracite is also an effective way to enhance the roasting performance of hematite pellet.

#### **4. Conclusions**

348 Recrystallization

Roasting time: a-4min; b-6min; c-8min; d-20min

(Anthracite: 1.0%, Oxygen partial pressure: 20%, airflow: 6L/min)

Fig. 17. Microstructure of the carbon-burdened pellets roasted at 1280C

1-Magnetite; 2-OH;3- SH

Both magnetite and burdened carbon are found to be the favourable techniques for enhancing the induration of hematite pellet. The induration mechanisms of hematite pellet with addition of magnetite concentrate and anthracite powder are revealed by characterization of Fe2O3 recrystallization rules during the oxidization roasting.

The crystallization behaviours of Fe2O3 during preheating and roasting of pellets made from mixed hematite/magnetite (H/M) concentrates have been revealed. The results indicate that the strength of pellet is mainly provided by the crystalline connections between the particles during preheating. This occurs because the activity of Fe2O3 from secondary hematite (oxidized from magnetite concentrate) is higher than that from original hematite (from the raw hematite concentrate). In the roasting process, when temperature is lower than 1250C, the strength is mainly provided by the development, connection and growth of Fe2O3 crystalline grains from secondary hematite. Only if the temperature exceeds 1280C does Fe2O3 recrystallization in original hematite grains develop very well.

**14**

*1,2,4Brazil 3France*

**Ion-Beam-Induced Epitaxial Recrystallization**

Shay Reboh3,4, Eliermes Meneses1, Livio Amaral4 and Lisandro Cardoso1

The transition from amorphous Silicon to crystalline Silicon is a process of great technological importance and has raised an enormous interest also from a purely scientific perspective. Ion irradiation through an amorphous/crystalline interface may stimulate recrystallization or layer-by-layer amorphization depending on the sample temperature and ion beam parameters. In this chapter, we address some key features of this recrystallization phenomenon. The recrystallization/amorphization process will be discussed in relation to its dependence on the energy deposited during the ion beam irradiation, the sample temperature, and the presence of impurities, such as iron atoms dissolved within an amorphous silicon layer. Also it will be discussed the specific experimental condition under which a metastable phase of the FeSi2 binary compound is trapped within a region of the recrystallized Silicon, in the form of nanoparticles. Theses nanoparticles, with different orientations and morphologies, are shown to cause interesting distortions in the surrounding Si crystal lattice. This effect reflects an important application of the ion-beam induced recrystallization process, as a method that could be used to synthesize ordered

The amorphous Silicon (a-Si) is a Si phase with well-defined thermodynamic properties and that presents Gibbs free energy of ~ 0.12 eV/at. higher than that of the crystalline phase (c-Si) (Donovan et al., 1985, 1989; Roorda et al., 1989). This implies the existence of a driving force for the transition from the amorphous to the crystalline phase to occur. In other words, there is a natural tendency for the (a-Si) → (c-Si) transformation. At room temperature, the a-Si phase is metastable and it is transformed into c-Si only when submitted to high temperatures, typically higher than 450 °C (Olson & Roth, 1988; Williams, 1983). For the case of an amorphized layer on top of a Si single-crystal substrate,

**1. Introduction**

nanoparticles within a Silicon matrix.

**2. Epitaxial crystallization**

**Method and Its Recent Applications**

*1Instituto de Física Gleb Wataghin - UNICAMP, Campinas, SP 2CCSST, Universidade Federal do Maranhão, Imperatriz, MA*

> *3Groupe nMat, CEMES-CNRS, Toulouse 4Instituto de Física - UFRGS, Porto Alegre, RS*

Rossano Lang1,4, Alan de Menezes1, Adenilson dos Santos2,

Because the activity of Fe2O3 in the secondary hematite grains is higher than that in the original hematite grains, Fe2O3 recrystallization bonds between particles can be enhanced by the secondary hematite grains during preheating and roasting, and the secondary hematite in H/M concentrate pellet is able to improve the strength and decrease the roasting temperature of hematite pellet. Thus, adding a certain proportion of magnetite concentrate is an effective way to improve the roasting performance of hematite pellet.

The effects of anthracite on oxidation roasting behaviour for hematite pellet have been elucidated. Anthracite in pellet has two functions: the one lies on that a part of heat needed in roasting process can be supplied by the carbon combustion, and the other is that the reduction of partial hematite by the carbon dispersed in pellet, as well as the partial decomposition of hematite at relatively low oxygen partial pressure, leads to the transformation of hematite into magnetite during the roasting.

Based on microstructure analysis, it can be founded that the new-born magnetite, produced from the reduction and the decomposition of original hematite (OH), is oxidized into the secondary hematite (SH) during roasting. Thus, Fe2O3 recrystallization bonds between particles can be consolidated by the secondary hematite grains at lower roasting temperature. The strength of carbon-burdened hematite pellet is enhanced and the roasting temperature is decreased due to the formation of secondary hematite. Therefore, adding a certain proportion of anthracite is also an effective way to enhance the roasting performance of hematite pellet.

#### **5. Acknowledgments**

The authors want to express their thanks to National Science Fund for Distinguished Young Scholars (50725416), National Natural Science Foundation of China (50604015 and 50804059) and Fundamental Research Funds for the Central Universities for financial supports of this research. Dr. Mingjun Rao is appreciated for his helpful remarks on spelling and expression.

#### **6. References**

### **Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications**

Rossano Lang1,4, Alan de Menezes1, Adenilson dos Santos2, Shay Reboh3,4, Eliermes Meneses1, Livio Amaral4 and Lisandro Cardoso1 *1Instituto de Física Gleb Wataghin - UNICAMP, Campinas, SP 2CCSST, Universidade Federal do Maranhão, Imperatriz, MA 3Groupe nMat, CEMES-CNRS, Toulouse 4Instituto de Física - UFRGS, Porto Alegre, RS 1,2,4Brazil 3France*

#### **1. Introduction**

350 Recrystallization

Because the activity of Fe2O3 in the secondary hematite grains is higher than that in the original hematite grains, Fe2O3 recrystallization bonds between particles can be enhanced by the secondary hematite grains during preheating and roasting, and the secondary hematite in H/M concentrate pellet is able to improve the strength and decrease the roasting temperature of hematite pellet. Thus, adding a certain proportion of magnetite concentrate

The effects of anthracite on oxidation roasting behaviour for hematite pellet have been elucidated. Anthracite in pellet has two functions: the one lies on that a part of heat needed in roasting process can be supplied by the carbon combustion, and the other is that the reduction of partial hematite by the carbon dispersed in pellet, as well as the partial decomposition of hematite at relatively low oxygen partial pressure, leads to the

Based on microstructure analysis, it can be founded that the new-born magnetite, produced from the reduction and the decomposition of original hematite (OH), is oxidized into the secondary hematite (SH) during roasting. Thus, Fe2O3 recrystallization bonds between particles can be consolidated by the secondary hematite grains at lower roasting temperature. The strength of carbon-burdened hematite pellet is enhanced and the roasting temperature is decreased due to the formation of secondary hematite. Therefore, adding a certain proportion of anthracite is also an effective way to enhance the roasting performance of hematite pellet.

The authors want to express their thanks to National Science Fund for Distinguished Young Scholars (50725416), National Natural Science Foundation of China (50604015 and 50804059) and Fundamental Research Funds for the Central Universities for financial supports of this research. Dr. Mingjun Rao is appreciated for his helpful remarks on spelling and expression.

Jiang, T.; Zhang, Y. & Huang, Z. (2008). Preheating and Roasting Characteristics of

Li, G.; Li, X. & Zhang, Y. (2009). Induration Mechanisms of Oxidised Pellets Prepared from

Xu, M. (2001). Deve1opment of BF Burden Structure in China. *Ironmaking*, Vol. 20, No. 2, pp.

Ball, D.; Dartnell, J. & Davison, J. (1973). *Agglomeration of Iron Ores*, Heinemann Educational,

APbill, J. (1985). Carbonaceous Additives for Pelletizing Production, *4TH International Symposium on Agglomeration*, ISBN 0-932897-00-2, Toronto, Canada, June, 1985 Clout, J.; Manuel J. (2003). Fundamental Investigations of Differences in Bonding

Yang, X.; Guo, Z. & Wang, D. (1995). Research on the Reduction Mechanism of Iron Ore

Ores. *Powder Technology*, Vol. 130, pp. (393-399), ISSN 0921-8831

Hematite–Magnetite (H–M) Concentrate Pellets. *Ironmaking Steelmaking*, Vol. 35,

Mixed Magnetite–Haematite Concentrates. *Ironmaking Steelmaking*, Vol. 36, No. 5,

Mechanisms in Iron Ore Sinter Formed from Magnetite Concentrates and Hematite

Pellets Containing Graphite. *Engineering Chemistry and Metallurgy*, Vol. 16, No. 2,

is an effective way to improve the roasting performance of hematite pellet.

transformation of hematite into magnetite during the roasting.

No. 1, pp. (21-26), ISSN 0301-9233

pp. (393-396), ISSN 0301-9233

(24-27), ISSN 1001-1471

ISBN<PHONE_NUMBER>, London

pp. (118-126), ISSN 1001-2052

**5. Acknowledgments**

**6. References**

The transition from amorphous Silicon to crystalline Silicon is a process of great technological importance and has raised an enormous interest also from a purely scientific perspective. Ion irradiation through an amorphous/crystalline interface may stimulate recrystallization or layer-by-layer amorphization depending on the sample temperature and ion beam parameters. In this chapter, we address some key features of this recrystallization phenomenon. The recrystallization/amorphization process will be discussed in relation to its dependence on the energy deposited during the ion beam irradiation, the sample temperature, and the presence of impurities, such as iron atoms dissolved within an amorphous silicon layer. Also it will be discussed the specific experimental condition under which a metastable phase of the FeSi2 binary compound is trapped within a region of the recrystallized Silicon, in the form of nanoparticles. Theses nanoparticles, with different orientations and morphologies, are shown to cause interesting distortions in the surrounding Si crystal lattice. This effect reflects an important application of the ion-beam induced recrystallization process, as a method that could be used to synthesize ordered nanoparticles within a Silicon matrix.

#### **2. Epitaxial crystallization**

The amorphous Silicon (a-Si) is a Si phase with well-defined thermodynamic properties and that presents Gibbs free energy of ~ 0.12 eV/at. higher than that of the crystalline phase (c-Si) (Donovan et al., 1985, 1989; Roorda et al., 1989). This implies the existence of a driving force for the transition from the amorphous to the crystalline phase to occur. In other words, there is a natural tendency for the (a-Si) → (c-Si) transformation. At room temperature, the a-Si phase is metastable and it is transformed into c-Si only when submitted to high temperatures, typically higher than 450 °C (Olson & Roth, 1988; Williams, 1983). For the case of an amorphized layer on top of a Si single-crystal substrate,

Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications 353

Rutherford backscattering spectrometry in combination with the ion channeling technique (RBS/C) is commonly used to directly monitor the c-a interface motion and infer the kinetics of the IBIEC process. The following RBS/C results were carried out with He+ beam at 1 MeV in a 170° scattering geometry. Figure 2 presents c-a interface position measurements of [100] oriented Si substrate previously implanted at room temperature with Fe+ ions at 100 keV energy. Channeled implantation was avoided by tilting the sample 7° normal with respect to the incident beam direction. Subsequently, the recrystallization of the Fe-implanted surface amorphous Si layer (filled circles in the figure) was induced by Si+ irradiation at high energy - 600 keV (whose projected range is well beyond the original c-a interface). In the IBIEC experiments the substrate temperature was fixed and controlled. This allows discrimination between the effects due to the heating of the sample holder and those due to ion-beam irradiation. In order to avoid beam heating effects low current density ( 1 μA/cm2) was therefore used with ion flux of ~ 6.2x1012 ions/cm2s. The substrate temperature was maintained at 350 °C. Essentially, one observes that the increase of the Si+ irradiated dose, leads to a decrease in the distance between the c-a interface and the surface. That is, there is

Fig. 2. Determination of the thickness of an amorphous layer using the ion channeling technique. This particular case corresponds to the recrystallization of ~ 176 nm a-Si under

In implantation processes (Dearnaley et al., 1973), the implanted dose Φ is the integral of the ion flux in time, which corresponds to the total number of ions that were focused on the

<sup>Φ</sup> <sup>Φ</sup> (ion/cm ) ; (s) A / cm

*e*

*nq*

0 2

*<sup>T</sup> <sup>e</sup>*

*<sup>I</sup> nq dt T*

, (1)

600 keV Si+ beam irradiation keeping the substrate at 350 °C.

2

sample per unit area

**2.1 Crystal-amorphous interface displacement velocity**

a recrystallization towards the surface.

the transition occurs by a planar motion of the crystal-amorphous (c-a) interface, from the interior towards the surface, and hence decreasing the thickness of the existing amorphous layer with time, as schematically illustrated in Fig. 1. This effect is called epitaxial crystallization. For a pure thermal recrystallization process - SPEG (*Solid Phase Epitaxial Growth*), the growth rate is strongly dependent on the temperature and it presents an Arrhenius-like behavior with a unique activation energy of (2.68 0.05) eV over a growth rate range of more than six orders of magnitude (Olson & Roth, 1988). For instance, at 470 °C the c-a interface displacement velocity is ~ 1 Å/min while at lower temperatures, this value decreases considerably and the amorphous to crystal transition becomes kinetically inhibited. However, epitaxial recrystallization of a-Si layers can also be achieved at lower temperatures (~ 200 - 320 °C) by ion-beam irradiation. This phenomenon represents a typical example of a dynamic annealing process and it is usually referred to as *Ion-Beam-Induced Epitaxial Crystallization* - IBIEC (Priolo & Rimini, 1990).

It is well known that both the implantation and the irradiation processes (depending on the dose and the beam ion mass) have the amorphization of the host matrix, as one of their main effects. For the specific case of Si, when the deposited energy (energy loss of ion beam mainly by nuclear collisions) handed by the projectile exceeds the threshold value of ~ 13 eV/at. a buried amorphous Si layer is formed (Narayan et al., 1984). However, if the Si is irradiated at a temperature above ~ 150 °C, there is a competition between both phenomena: amorphization and recrystallization.

Fig. 1. Schematic illustration of two recrystallization processes in Silicon. SPEG (Solid Phase Epitaxial Growth) is a purely thermal process while, IBIEC is an ion beam induced epitaxial recrystallization process.

In fact, the temperature rise of the Si substrate is not solely responsible for this trend of the crystalline order recovery, since the recrystallization rate is greater than that obtained by a purely thermal process at the same temperature. A more complex mechanism involving a dynamic reordering stimulated by the ion beam is responsible for this effect. The recrystallization under ion-beam irradiation is the result of the prevalence of dynamic annealing over defect generation rate. Both processes depend on the substrate temperature and on the parameters of the irradiating beam, such as ion species, energy, ion dose and ion flux - also known as dose rate.

#### **2.1 Crystal-amorphous interface displacement velocity**

352 Recrystallization

the transition occurs by a planar motion of the crystal-amorphous (c-a) interface, from the interior towards the surface, and hence decreasing the thickness of the existing amorphous layer with time, as schematically illustrated in Fig. 1. This effect is called epitaxial crystallization. For a pure thermal recrystallization process - SPEG (*Solid Phase Epitaxial Growth*), the growth rate is strongly dependent on the temperature and it presents an Arrhenius-like behavior with a unique activation energy of (2.68 0.05) eV over a growth rate range of more than six orders of magnitude (Olson & Roth, 1988). For instance, at 470 °C the c-a interface displacement velocity is ~ 1 Å/min while at lower temperatures, this value decreases considerably and the amorphous to crystal transition becomes kinetically inhibited. However, epitaxial recrystallization of a-Si layers can also be achieved at lower temperatures (~ 200 - 320 °C) by ion-beam irradiation. This phenomenon represents a typical example of a dynamic annealing process and it is usually referred to as *Ion-Beam-Induced Epitaxial Crystallization* - IBIEC (Priolo & Rimini,

It is well known that both the implantation and the irradiation processes (depending on the dose and the beam ion mass) have the amorphization of the host matrix, as one of their main effects. For the specific case of Si, when the deposited energy (energy loss of ion beam mainly by nuclear collisions) handed by the projectile exceeds the threshold value of ~ 13 eV/at. a buried amorphous Si layer is formed (Narayan et al., 1984). However, if the Si is irradiated at a temperature above ~ 150 °C, there is a competition between both phenomena:

Fig. 1. Schematic illustration of two recrystallization processes in Silicon. SPEG (Solid Phase Epitaxial Growth) is a purely thermal process while, IBIEC is an ion beam induced epitaxial

In fact, the temperature rise of the Si substrate is not solely responsible for this trend of the crystalline order recovery, since the recrystallization rate is greater than that obtained by a purely thermal process at the same temperature. A more complex mechanism involving a dynamic reordering stimulated by the ion beam is responsible for this effect. The recrystallization under ion-beam irradiation is the result of the prevalence of dynamic annealing over defect generation rate. Both processes depend on the substrate temperature and on the parameters of the irradiating beam, such as ion species, energy, ion dose and ion

1990).

amorphization and recrystallization.

recrystallization process.

flux - also known as dose rate.

Rutherford backscattering spectrometry in combination with the ion channeling technique (RBS/C) is commonly used to directly monitor the c-a interface motion and infer the kinetics of the IBIEC process. The following RBS/C results were carried out with He+ beam at 1 MeV in a 170° scattering geometry. Figure 2 presents c-a interface position measurements of [100] oriented Si substrate previously implanted at room temperature with Fe+ ions at 100 keV energy. Channeled implantation was avoided by tilting the sample 7° normal with respect to the incident beam direction. Subsequently, the recrystallization of the Fe-implanted surface amorphous Si layer (filled circles in the figure) was induced by Si+ irradiation at high energy - 600 keV (whose projected range is well beyond the original c-a interface). In the IBIEC experiments the substrate temperature was fixed and controlled. This allows discrimination between the effects due to the heating of the sample holder and those due to ion-beam irradiation. In order to avoid beam heating effects low current density ( 1 μA/cm2) was therefore used with ion flux of ~ 6.2x1012 ions/cm2s. The substrate temperature was maintained at 350 °C. Essentially, one observes that the increase of the Si+ irradiated dose, leads to a decrease in the distance between the c-a interface and the surface. That is, there is a recrystallization towards the surface.

Fig. 2. Determination of the thickness of an amorphous layer using the ion channeling technique. This particular case corresponds to the recrystallization of ~ 176 nm a-Si under 600 keV Si+ beam irradiation keeping the substrate at 350 °C.

In implantation processes (Dearnaley et al., 1973), the implanted dose Φ is the integral of the ion flux in time, which corresponds to the total number of ions that were focused on the sample per unit area

$$\Phi \text{ (ion/cm}^2\text{)} = \int\_0^T \frac{I}{n\eta\_e} dt \qquad ; \quad T(\text{s}) = \frac{\Phi \cdot n\eta\_e}{\text{A / cm}^2} , \tag{1}$$

Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications 355

determined as shown in figure 4. In the same figure, one observes the number of vacancies produced by irradiation of Si+ ions as a function of depth. This evaluation corresponds to a calculation obtained by the SRIM algorithm with a displacement energy of 15 eV and a lattice binding energy of 2 eV) (Ziegler, 2011). The depth dependence of the regrowth

Fig. 4. Regrowth rate versus depth for a Si layer amorphized by Fe+ ions and recrystallized at 350 °C by 600 keV Si+ ions. The inset shows the number of vacancies generated by Si

This result suggests that the ion-induced recrystallization rate is associated with the production of point defects, or in a more general way, with the energy loss into elastic collisions at the c-a interface. Therefore, only those defects generated nearby or directly at the c-a interface are available for the recrystallization process. However, it should be mentioned that the diffusion or vacancy migration is not the promoter mechanism of the IBIEC. This issue will be addressed later. The abovementioned conclusions were also confirmed by other experimental results (Holmén et al., 1984; Linnros et al., 1984, 1985; Miyao et al., 1986; Williams et al., 1985). For instance, Linnros (Linnros et al., 1985) reported a clear experimental evidence of a linear dependence of the ion-induced recrystallization rate on the nuclear energy loss. In their experiments, ion beams of different masses (He, N, Ne, Si, As, Kr) were used to stimulate IBIEC where the recrystallization rate was observed to increase with increasing ion mass. This demonstrates that nuclear energy loss is the mechanism which produces the defects responsible for the IBIEC. Furthermore, the effects produced by both kinds of energy loss (nuclear and electronic) were discriminated through the dependence of the recrystallization rate on the beam energy. It was inferred that the electronic excitations and ionizations play practically no role in the recrystallization process (Elliman et al., 1985; Williams et al., 1985). In other experiments, in which a-Si layers were irradiated by the electron-beam, the recrystallization was observed only for energies above a threshold for atomic displacement of ~ 150 keV while below this threshold no epitaxial

experimental rates seems to follow the profile of defect generation.

irradiation calculated by the SRIM code (closed circles).

where: *I* is the beam current in ampere per unit area of the sample (A/cm2), *qe* is the electron charge, *n* = 1 for once ionized ions, *n* = 2 for doubly ionized species, and so on. The term *I*/*qe* is designated as the ion flux (ion/cm2s) and *T*(s) is the implantation time in seconds. For the IBIEC, one designates a certain irradiation dose Φ and, consequently, the ion flux for such irradiation by = *d*Φ/*dt*. When the ion flux is kept constant at a given crystallization procedure, the irradiation dose may be interpreted as a measure of the processing time described by

$$\frac{d\mathbf{X}\_{c-a}}{dt} = \frac{d\mathbf{X}\_{c-a}}{d\boldsymbol{\Phi}} \frac{d\boldsymbol{\Phi}}{dt} = \boldsymbol{\phi} \frac{d\mathbf{X}\_{c-a}}{d\boldsymbol{\Phi}} \,. \tag{2}$$

Several reports define a c-a interface velocity *R* as the derivative of the recrystallized thickness (nm) *versus* dose (at/cm2) curve. Thus, the unit of such "velocity" is expressed in nm/(at/cm2). In the present study, figure 3 is an example that enables us to extract the interface displacement velocity or recrystallization rate.

Fig. 3. Regrown thickness for [100] Si amorphized by Fe+ ions and subsequently irradiated at 350 °C by 600 keV Si+ ions. The interfacial displacement velocity *R* can be extracted from the curve.

It should be noted that a comparison between the thermal crystallization rate *V* ≡ (d*Xc*-*<sup>a</sup>*/d*t*) and the recrystallization rate defined for the IBIEC *R* ≡ (*dXc*-*<sup>a</sup>*/*d*Φ) is always possible, if one knows the ion flux of the process as shown in equation (2).

#### **2.2 Influence of the beam parameters in the IBIEC process**

Several experimental results have indicated that the beam parameters such as ion species (by the factor of nuclear energy loss *Sn*) and the ion flux have direct influence on the recrystallization rate. With regards to the energy loss by projectile-target elastic collisions, the following result is quite interesting. By using the recrystallized thickness *versus* irradiation dose plot (Fig. 3), the recrystallization rate *R* as a function of depth was

where: *I* is the beam current in ampere per unit area of the sample (A/cm2), *qe* is the electron charge, *n* = 1 for once ionized ions, *n* = 2 for doubly ionized species, and so on. The term *I*/*qe*

the IBIEC, one designates a certain irradiation dose Φ and, consequently, the ion flux for

procedure, the irradiation dose may be interpreted as a measure of the processing time

d Φ Φ *X dX dX ca ca c a d t d dt d*

Several reports define a c-a interface velocity *R* as the derivative of the recrystallized thickness (nm) *versus* dose (at/cm2) curve. Thus, the unit of such "velocity" is expressed in nm/(at/cm2). In the present study, figure 3 is an example that enables us to extract the

Fig. 3. Regrown thickness for [100] Si amorphized by Fe+ ions and subsequently irradiated at 350 °C by 600 keV Si+ ions. The interfacial displacement velocity *R* can be extracted from the

It should be noted that a comparison between the thermal crystallization rate *V* ≡ (d*Xc*-*<sup>a</sup>*/d*t*) and the recrystallization rate defined for the IBIEC *R* ≡ (*dXc*-*<sup>a</sup>*/*d*Φ) is always possible, if one

Several experimental results have indicated that the beam parameters such as ion species

recrystallization rate. With regards to the energy loss by projectile-target elastic collisions, the following result is quite interesting. By using the recrystallized thickness *versus* irradiation dose plot (Fig. 3), the recrystallization rate *R* as a function of depth was

have direct influence on the

of the process as shown in equation (2).

**2.2 Influence of the beam parameters in the IBIEC process**

(by the factor of nuclear energy loss *Sn*) and the ion flux

d Φ

(ion/cm2s) and *T*(s) is the implantation time in seconds. For

. (2)

= *d*Φ/*dt*. When the ion flux is kept constant at a given crystallization

is designated as the ion flux

such irradiation by

described by

curve.

knows the ion flux

interface displacement velocity or recrystallization rate.

determined as shown in figure 4. In the same figure, one observes the number of vacancies produced by irradiation of Si+ ions as a function of depth. This evaluation corresponds to a calculation obtained by the SRIM algorithm with a displacement energy of 15 eV and a lattice binding energy of 2 eV) (Ziegler, 2011). The depth dependence of the regrowth experimental rates seems to follow the profile of defect generation.

Fig. 4. Regrowth rate versus depth for a Si layer amorphized by Fe+ ions and recrystallized at 350 °C by 600 keV Si+ ions. The inset shows the number of vacancies generated by Si irradiation calculated by the SRIM code (closed circles).

This result suggests that the ion-induced recrystallization rate is associated with the production of point defects, or in a more general way, with the energy loss into elastic collisions at the c-a interface. Therefore, only those defects generated nearby or directly at the c-a interface are available for the recrystallization process. However, it should be mentioned that the diffusion or vacancy migration is not the promoter mechanism of the IBIEC. This issue will be addressed later. The abovementioned conclusions were also confirmed by other experimental results (Holmén et al., 1984; Linnros et al., 1984, 1985; Miyao et al., 1986; Williams et al., 1985). For instance, Linnros (Linnros et al., 1985) reported a clear experimental evidence of a linear dependence of the ion-induced recrystallization rate on the nuclear energy loss. In their experiments, ion beams of different masses (He, N, Ne, Si, As, Kr) were used to stimulate IBIEC where the recrystallization rate was observed to increase with increasing ion mass. This demonstrates that nuclear energy loss is the mechanism which produces the defects responsible for the IBIEC. Furthermore, the effects produced by both kinds of energy loss (nuclear and electronic) were discriminated through the dependence of the recrystallization rate on the beam energy. It was inferred that the electronic excitations and ionizations play practically no role in the recrystallization process (Elliman et al., 1985; Williams et al., 1985). In other experiments, in which a-Si layers were irradiated by the electron-beam, the recrystallization was observed only for energies above a threshold for atomic displacement of ~ 150 keV while below this threshold no epitaxial

Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications 357

It should be noted that recrystallization due to ion-beam irradiation occurs in a temperature range for which the thermal process is kinetically inhibited. Therefore, ion-beam irradiation strongly enhances the kinetics of recrystallization. For instance, at 250 °C the ion-induced growth rate is 0.07 Å/s while, an extrapolation of the thermal data gives a rate of only 10-10 Å/s. Furthermore, in the temperature range shown, the IBIEC presents an Arrhenius-like temperature dependence with an apparent activation energy of (0.32 0.05) eV as also demonstrated by several other experiments (Elliman et al., 1985; Linnros et al., 1984;

It has been proposed that the activation energy for thermal recrystallization of a-Si layers is composed of two terms: one for defect generation and other for defect migration. Ion-beam irradiation clearly removes the main activated process usually limiting conventional thermal regrowth. In fact, during ion-beam irradiation defects are not thermally generated but rather being produced by means of atomic collisions. Jackson (Jackson, 1988) proposed that the activation energy inferred from the IBIEC is not associated with any activated process and therefore considered just as an apparent activation energy. The linear dependence of the regrowth rate as a function of reciprocal temperature, in a logarithmic plot, comes instead, from a balance between different effects. In addition, Jackson (Jackson, 1988) in his intracascade model suggested the dangling bond in the amorphous phase as the promoter for IBIEC. Dangling bonds are structural defects which by moving nearby the c-a interfacial region should produce a rearrangement of the bonds leading to recrystallization. In the mid-1990s, Priolo (Priolo et al., 1990) proposed a phenomenological model of IBIEC which combines the approach of the Jackson model (Jackson, 1988) with the structural and electronic features of models proposed for conventional thermal regrowth (Williams &

Elliman, 1986). This model has explained all the experimental results so far.

In the previous sections, it was shown how a combination of thermal energy and energy deposited by ballistic effects can produce a non-equilibrium epitaxial recrystallization. However, IBIEC is a reversible process, where the increase in the ion flux of the irradiating beam, and /or the decrease in the target temperature can cause a planar layer-by-layer amorphization instead of an epitaxial recrystallization (Elliman et al., 1987; Linnros et al., 1986, 1988). Both processes are schematically illustrated in figure 6. At a constant flux, there is a substrate temperature TR (reversal temperature) such that, when T > TR, the ion irradiation produces epitaxial regrowth whereas, when T < TR the irradiation produces a layer-by-layer amorphization. The remarkable point is that the amorphization occurs just from the pre-existing amorphous seed and not from regions below the c-a interface, despite the fact that the energy loss by elastic collisions increases with the increasing depth. In this context, the net velocity of the interface motion can be described by two terms: a

d dd / / <sup>Φ</sup> d dd

When *R* is positive, the crystallization regime is prevalent and when *R* is negative the

. (3)

*ca ca c a dX X X X <sup>R</sup> d t tt*

Williams et al., 1985).

**2.4 Planar amorphization**

crystallization and an amorphization, as

system is in the amorphization regime.

regrowth was produced, even after irradiation at very high electron doses (Lulli et al., 1987; Miyao et al., 1986; Washburn et al., 1983). Therefore, the observed epitaxy is associated with elastic collisions that transfer sufficient momentum to displace target atoms from their lattice site.

As mentioned earlier in this section, there is another parameter that directly influences the process - the ion flux. In general, it is observed that the lower the flux , the higher is the effective velocity *R* ≡ (*dXc*-*<sup>a</sup>*/*d*Φ) defined by equation 2. This parameter is so important that high fluxes almost inhibit the process, especially for heavier ions (Linnros et al., 1985).

#### **2.3 Temperature dependence**

The sample temperature during irradiation is a fundamental variable in the IBIEC process (Elliman et al., 1985; Linnros et al., 1984; Priolo et al., 1988, 1989; Williams et al., 1985). Figure 5 shows, in an Arrhenius plot, the ion-induced growth rate (or recrystallization) as a function of the reciprocal temperature of an a-Si/Si(100) layer recrystallized by 600 keV Kr2+ ion irradiation (dose: 1x1015 Kr/cm2 and ion flux: 1x1012 Kr/cm2s). It also shows the recrystallization rate that represents the thermal contribution (SPEG) with an activation energy of (2.68 0.05) eV. The data were extracted from references (Olson & Roth, 1988; Priolo et al., 1988) and reported in the figure. The growth rate is reported in Å/s (left-hand side) and in Å4/eV (right-hand side). The latter scale represents the growth rate in the form ΔX/Φ*ν*(E), ΔX being the recrystallized thickness, Φ the dose and *ν*(E) the total energy deposited responsible for the displacement production at the c-a interface.

Fig. 5. Ion-induced growth rate versus reciprocal temperature for a-Si layers produced by Si+ implantation and recrystallized by 600 keV Kr2+ ion irradiation. The thermal contribution to the growth rate is also displayed. Data extracted from references (Olson & Roth, 1988; Priolo et al., 1988) and reported in the figure.

It should be noted that recrystallization due to ion-beam irradiation occurs in a temperature range for which the thermal process is kinetically inhibited. Therefore, ion-beam irradiation strongly enhances the kinetics of recrystallization. For instance, at 250 °C the ion-induced growth rate is 0.07 Å/s while, an extrapolation of the thermal data gives a rate of only 10-10 Å/s. Furthermore, in the temperature range shown, the IBIEC presents an Arrhenius-like temperature dependence with an apparent activation energy of (0.32 0.05) eV as also demonstrated by several other experiments (Elliman et al., 1985; Linnros et al., 1984; Williams et al., 1985).

It has been proposed that the activation energy for thermal recrystallization of a-Si layers is composed of two terms: one for defect generation and other for defect migration. Ion-beam irradiation clearly removes the main activated process usually limiting conventional thermal regrowth. In fact, during ion-beam irradiation defects are not thermally generated but rather being produced by means of atomic collisions. Jackson (Jackson, 1988) proposed that the activation energy inferred from the IBIEC is not associated with any activated process and therefore considered just as an apparent activation energy. The linear dependence of the regrowth rate as a function of reciprocal temperature, in a logarithmic plot, comes instead, from a balance between different effects. In addition, Jackson (Jackson, 1988) in his intracascade model suggested the dangling bond in the amorphous phase as the promoter for IBIEC. Dangling bonds are structural defects which by moving nearby the c-a interfacial region should produce a rearrangement of the bonds leading to recrystallization. In the mid-1990s, Priolo (Priolo et al., 1990) proposed a phenomenological model of IBIEC which combines the approach of the Jackson model (Jackson, 1988) with the structural and electronic features of models proposed for conventional thermal regrowth (Williams & Elliman, 1986). This model has explained all the experimental results so far.

#### **2.4 Planar amorphization**

356 Recrystallization

regrowth was produced, even after irradiation at very high electron doses (Lulli et al., 1987; Miyao et al., 1986; Washburn et al., 1983). Therefore, the observed epitaxy is associated with elastic collisions that transfer sufficient momentum to displace target atoms from their

As mentioned earlier in this section, there is another parameter that directly influences the

effective velocity *R* ≡ (*dXc*-*<sup>a</sup>*/*d*Φ) defined by equation 2. This parameter is so important that high fluxes almost inhibit the process, especially for heavier ions (Linnros et al., 1985).

The sample temperature during irradiation is a fundamental variable in the IBIEC process (Elliman et al., 1985; Linnros et al., 1984; Priolo et al., 1988, 1989; Williams et al., 1985). Figure 5 shows, in an Arrhenius plot, the ion-induced growth rate (or recrystallization) as a function of the reciprocal temperature of an a-Si/Si(100) layer recrystallized by 600 keV Kr2+ ion irradiation (dose: 1x1015 Kr/cm2 and ion flux: 1x1012 Kr/cm2s). It also shows the recrystallization rate that represents the thermal contribution (SPEG) with an activation energy of (2.68 0.05) eV. The data were extracted from references (Olson & Roth, 1988; Priolo et al., 1988) and reported in the figure. The growth rate is reported in Å/s (left-hand side) and in Å4/eV (right-hand side). The latter scale represents the growth rate in the form ΔX/Φ*ν*(E), ΔX being the recrystallized thickness, Φ the dose and *ν*(E) the total energy

Fig. 5. Ion-induced growth rate versus reciprocal temperature for a-Si layers produced by Si+ implantation and recrystallized by 600 keV Kr2+ ion irradiation. The thermal contribution to the growth rate is also displayed. Data extracted from references (Olson & Roth, 1988; Priolo

, the higher is the

process - the ion flux. In general, it is observed that the lower the flux

deposited responsible for the displacement production at the c-a interface.

lattice site.

**2.3 Temperature dependence**

et al., 1988) and reported in the figure.

In the previous sections, it was shown how a combination of thermal energy and energy deposited by ballistic effects can produce a non-equilibrium epitaxial recrystallization. However, IBIEC is a reversible process, where the increase in the ion flux of the irradiating beam, and /or the decrease in the target temperature can cause a planar layer-by-layer amorphization instead of an epitaxial recrystallization (Elliman et al., 1987; Linnros et al., 1986, 1988). Both processes are schematically illustrated in figure 6. At a constant flux, there is a substrate temperature TR (reversal temperature) such that, when T > TR, the ion irradiation produces epitaxial regrowth whereas, when T < TR the irradiation produces a layer-by-layer amorphization. The remarkable point is that the amorphization occurs just from the pre-existing amorphous seed and not from regions below the c-a interface, despite the fact that the energy loss by elastic collisions increases with the increasing depth. In this context, the net velocity of the interface motion can be described by two terms: a crystallization and an amorphization, as

$$R = \frac{dX\_{c-a}}{d\Phi} = \frac{dX\_{c-a}}{\mathbf{d}t} / \left\langle \phi = \left[\frac{\mathbf{dX}\_c}{\mathbf{d}t} - \frac{\mathbf{dX}\_a}{\mathbf{d}t}\right] / \phi \right. \tag{3}$$

When *R* is positive, the crystallization regime is prevalent and when *R* is negative the system is in the amorphization regime.

Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications 359

used during the IBIEC experiments, since the interface displacement velocity is much higher than the values of their mobilities in amorphous silicon. In this case, the initial concentration profile of these impurities remains unchanged after recrystallization. This allows one to produce non-equilibrium structures with impurities trapped in the c-Si at concentrations well above their solid solubility. However, the presence of impurities at the c-a interface can modify the recrystallization rate which is observed to increase or decrease, according to the

Fig. 7. RBS channeling spectra along the [100] Si axis for a ~ 90 nm thick amorphous layer before (open circles) and after (closed circles) 380 keV Ne+ irradiation at 100 °C and dose of

From this brief discussion, it is noticed that the behavior of crucial parameters in the IBIEC process such as temperature and impurities is very similar to that in solid-phase epitaxial growth case. Despite the fact that the impact of these parameters is quite different for each phenomenon, there is a strong indication that similar microscopic processes occur in both cases. More precisely, the same interfacial defect responsible for the thermal recrystallization is considered to be active in IBIEC, being the ion beam the precursor of the increase in the

The IBIEC technique has been used as a method to synthesize nanoparticles in Silicon matrix [Lang et al., 2010a, 2010b]. Specifically, the structural properties of the FeSi2 nanoparticles synthesized in Fe+ low dose implanted Si(100) substrates were investigated. Particularly in this experimental observation, the Fe proved to be a fast diffuser for IBIEC, despite of its action as a retardant of the process, whose recrystallization rate was dependent on the implanted Fe+ concentration. Nevertheless, only the main results are reported here. The remarkable results which are presented show that the shape of the synthesized material

1x1017 /cm2. Note the planar growth of the amorphous phase.

average concentration of these defects.

**3. Nanoparticles synthesis by IBIEC**

particular species and to its concentration.

Fig. 6. Schematic representation of the ion-induced crystallization and amorphization. At temperatures below TR the amorphous layer grows (amorphization regime), whereas at temperatures above TR, it shrinks (crystallization regime).

An example of planar amorphization is shown in figure 7. In this example, a ~ 90 nm thick layer onto a Si(100) substrate was amorphized at room temperature by 40 keV Fe+ ions at a dose of 1x1016 /cm2. The irradiation was performed by using a 380 KeV Ne+ beam at a dose of 1x1017 /cm2 and ion flux of ~ 1.5x1013/cm2s. The substrate temperature was fixed at 100 °C. In the figure two RBS spectra are displayed in channeling condition for the a-Si layer before (open circles) and after (closed circles) Ne irradiation. A random spectrum is also reported. Ne irradiation clearly produces a great amount of damage beyond the original c-a interface. The a-Si surface layer is clearly seen to enlarge under a planar motion towards the sample interior. The amorphous layer has become ~ 50 nm thicker.

#### **2.5 Impurity effects**

Besides temperature and beam parameters, the kind of impurity dissolved in the sample plays an important role in the IBIEC. The presence of impurities dispersed within the a-Si layer can dramatically affect the recrystallization process. Depending on their behavior, they can be divided in two major categories: fast and slow diffusers. Fast diffusers comprehend species like Cu, Ag and Au which, at typical temperatures of the IBIEC process (~ 300 °C) have diffusivities of the order of 10-12 - 10-15 cm2/s and low solid solubility. These impurities have therefore enough mobility to be redistributed at the advancing c-a interface during recrystallization, modifying the impurity initial profile through the segregation towards the surface imposed by the planar advance. On the other hand, the slow diffusers such as B, P, As, do not present the effect of segregation - are immobile in the time-temperature windows

Fig. 6. Schematic representation of the ion-induced crystallization and amorphization. At temperatures below TR the amorphous layer grows (amorphization regime), whereas at

An example of planar amorphization is shown in figure 7. In this example, a ~ 90 nm thick layer onto a Si(100) substrate was amorphized at room temperature by 40 keV Fe+ ions at a dose of 1x1016 /cm2. The irradiation was performed by using a 380 KeV Ne+ beam at a dose of 1x1017 /cm2 and ion flux of ~ 1.5x1013/cm2s. The substrate temperature was fixed at 100 °C. In the figure two RBS spectra are displayed in channeling condition for the a-Si layer before (open circles) and after (closed circles) Ne irradiation. A random spectrum is also reported. Ne irradiation clearly produces a great amount of damage beyond the original c-a interface. The a-Si surface layer is clearly seen to enlarge under a planar motion towards the

Besides temperature and beam parameters, the kind of impurity dissolved in the sample plays an important role in the IBIEC. The presence of impurities dispersed within the a-Si layer can dramatically affect the recrystallization process. Depending on their behavior, they can be divided in two major categories: fast and slow diffusers. Fast diffusers comprehend species like Cu, Ag and Au which, at typical temperatures of the IBIEC process (~ 300 °C) have diffusivities of the order of 10-12 - 10-15 cm2/s and low solid solubility. These impurities have therefore enough mobility to be redistributed at the advancing c-a interface during recrystallization, modifying the impurity initial profile through the segregation towards the surface imposed by the planar advance. On the other hand, the slow diffusers such as B, P, As, do not present the effect of segregation - are immobile in the time-temperature windows

temperatures above TR, it shrinks (crystallization regime).

sample interior. The amorphous layer has become ~ 50 nm thicker.

**2.5 Impurity effects**

used during the IBIEC experiments, since the interface displacement velocity is much higher than the values of their mobilities in amorphous silicon. In this case, the initial concentration profile of these impurities remains unchanged after recrystallization. This allows one to produce non-equilibrium structures with impurities trapped in the c-Si at concentrations well above their solid solubility. However, the presence of impurities at the c-a interface can modify the recrystallization rate which is observed to increase or decrease, according to the particular species and to its concentration.

Fig. 7. RBS channeling spectra along the [100] Si axis for a ~ 90 nm thick amorphous layer before (open circles) and after (closed circles) 380 keV Ne+ irradiation at 100 °C and dose of 1x1017 /cm2. Note the planar growth of the amorphous phase.

From this brief discussion, it is noticed that the behavior of crucial parameters in the IBIEC process such as temperature and impurities is very similar to that in solid-phase epitaxial growth case. Despite the fact that the impact of these parameters is quite different for each phenomenon, there is a strong indication that similar microscopic processes occur in both cases. More precisely, the same interfacial defect responsible for the thermal recrystallization is considered to be active in IBIEC, being the ion beam the precursor of the increase in the average concentration of these defects.

#### **3. Nanoparticles synthesis by IBIEC**

The IBIEC technique has been used as a method to synthesize nanoparticles in Silicon matrix [Lang et al., 2010a, 2010b]. Specifically, the structural properties of the FeSi2 nanoparticles synthesized in Fe+ low dose implanted Si(100) substrates were investigated. Particularly in this experimental observation, the Fe proved to be a fast diffuser for IBIEC, despite of its action as a retardant of the process, whose recrystallization rate was dependent on the implanted Fe+ concentration. Nevertheless, only the main results are reported here. The remarkable results which are presented show that the shape of the synthesized material

Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications 361

the two distinct regions (crystalline and amorphous) is mainly composed of dangling bonds, and this particular kind of structural defect is responsible for the IBIEC process. Upon IBIEC conditions (temperature + irradiation) there is a dynamic rearrangement of these dangling bonds with annihilation in pairs which promotes the layer-by-layer planar recrystallization

Fig. 8. 1 MeV He+ RBS spectra in channeling along the [100] direction from a 90 nm a-Si layer containing Fe before (as-implanted sample) and after (recrystallized sample)

Fig. 9. Bright-field TEM cross-sectional images of the as-implanted sample. **(a)** Overview of the 90 nm thick a-Si layer containing Fe. **(b)** High-resolution of the selected region in **(a)**

irradiation with 600 keV Si+ to a dose of 6x1016 ions/cm2 at 350 °C.

showing in detail the c-a interface region.

toward the surface.

(observed by high resolution transmission electron microscopy - HRTEM) singularly affects the surrounding Si lattice. The lattice strain shape-dependent distribution in both directions: out-of-plane and in-plane was tailored by X-ray Bragg-Surface Diffraction technique.

#### **3.1 Fe<sup>+</sup> ion implantation in Si(100) and recrystallization process**

A Si(100) *n*-type Czochralski wafer (thickness 500 μm, resistivity 10 - 20 Ω cm) was used as host matrix. Fe+ ions at 40 keV were implanted at room temperature at an ion dose of 5x1015 cm-2. Channeling effects were avoided by tilting the sample 7º normal with respect to the incident beam direction. The typical iron beam current density during implantation was about 150 nA/cm2. Subsequently, ion-beam recrystallization experiments were performed at 350 °C using a 600 keV Si+ beam (current density 1 μA/cm2) to a total dose of 6x1016 ions/cm2. The dose rate resulting from the Si beam current was 6.2x1012 ions/cm2s. It is worth noting that Si+ ions at 600 keV energy have a projected range (average depth) of ~ 770 nm with a straggle (standard deviation) of ~ 150 nm. These values ensure that the irradiation exceeds the pre-existing amorphous layer (well beyond the original c-a interface).

The structures obtained into the as-implanted and recrystallized samples were analyzed and characterized by Rutherford Backscattering Spectrometry combined with ion channeling technique (RBS/C - with He+ beam at 1 MeV in a 170° scattering geometry) and also by transmission electron microscopy (TEM - JEOL 2010 operating at 200 kV). High-resolution rocking curves, as well as the reflection mappings of the Bragg-Surface Diffraction reflections, i.e., at the exact multiple diffraction condition, were carried out using the Huber multiaxis diffractometer of the XRD1 beam-line (Brazilian Synchrotron Radiation Facility - LNLS), with an incident beam wavelength of = 1.5495(5) Å, as defined by using a Si(111) channel-cut monochromator.

#### **3.1.1 RBS and TEM – Results and discussion**

Figure 8 shows RBS/C spectra obtained at random and [100]-channeled direction from the samples before (as-implanted) and after (recrystallized) irradiation. As observed in the aligned as-implanted spectrum, the implantation has produced an amorphous layer over ~ 90 nm while no channeling in the Fe signal was observed. However, expressive reduction of dechanneling yield in the Si profile is detected after IBIEC process (aligned IBIEC spectra relative to aligned as-implanted spectra). This decrease reflects the crystalline order recovery. The entire amorphous Si layer was recrystallized (at an average rate of ~ 0.04 nm/s) with a minimum backscattering yield on the subsurface region being χmin ≈ 6.3 %. The RBS concentration-depth profile (not shown here) has indicated an implanted Fe peak concentration of ~ 2.8 at.% at 40 nm from the surface. However, the recrystallization process caused a slight narrowing of the Fe peak and a small segregation towards the surface. A significant degree of channeling (χFe ≈ 46 %) was also observed in the Fe spectrum.

Figure 9 shows bright-field TEM cross-section images (taken at [110]Si zone axis) of the asimplanted sample. One clearly observes (Fig. 9a) the amorphous Si layer of ~ 90 nm produced by 40 keV Fe+ implantation and the crystal-amorphous interface. Figure 9b shows in detail the c-a interface, where one can note a large amount of defects called "end-of-range defects" generated by the implantation process. This defective intermediate zone between

(observed by high resolution transmission electron microscopy - HRTEM) singularly affects the surrounding Si lattice. The lattice strain shape-dependent distribution in both directions:

A Si(100) *n*-type Czochralski wafer (thickness 500 μm, resistivity 10 - 20 Ω cm) was used as host matrix. Fe+ ions at 40 keV were implanted at room temperature at an ion dose of 5x1015 cm-2. Channeling effects were avoided by tilting the sample 7º normal with respect to the incident beam direction. The typical iron beam current density during implantation was about 150 nA/cm2. Subsequently, ion-beam recrystallization experiments were performed at 350 °C using a 600 keV Si+ beam (current density 1 μA/cm2) to a total dose of 6x1016 ions/cm2. The dose rate resulting from the Si beam current was 6.2x1012 ions/cm2s. It is worth noting that Si+ ions at 600 keV energy have a projected range (average depth) of ~ 770 nm with a straggle (standard deviation) of ~ 150 nm. These values ensure that the irradiation exceeds the pre-existing amorphous layer (well beyond the original c-a

The structures obtained into the as-implanted and recrystallized samples were analyzed and characterized by Rutherford Backscattering Spectrometry combined with ion channeling technique (RBS/C - with He+ beam at 1 MeV in a 170° scattering geometry) and also by transmission electron microscopy (TEM - JEOL 2010 operating at 200 kV). High-resolution rocking curves, as well as the reflection mappings of the Bragg-Surface Diffraction reflections, i.e., at the exact multiple diffraction condition, were carried out using the Huber multiaxis diffractometer of the XRD1 beam-line (Brazilian Synchrotron Radiation Facility - LNLS), with an incident beam wavelength of = 1.5495(5) Å, as defined by using a Si(111)

Figure 8 shows RBS/C spectra obtained at random and [100]-channeled direction from the samples before (as-implanted) and after (recrystallized) irradiation. As observed in the aligned as-implanted spectrum, the implantation has produced an amorphous layer over ~ 90 nm while no channeling in the Fe signal was observed. However, expressive reduction of dechanneling yield in the Si profile is detected after IBIEC process (aligned IBIEC spectra relative to aligned as-implanted spectra). This decrease reflects the crystalline order recovery. The entire amorphous Si layer was recrystallized (at an average rate of ~ 0.04 nm/s) with a minimum backscattering yield on the subsurface region being χmin ≈ 6.3 %. The RBS concentration-depth profile (not shown here) has indicated an implanted Fe peak concentration of ~ 2.8 at.% at 40 nm from the surface. However, the recrystallization process caused a slight narrowing of the Fe peak and a small segregation towards the surface. A

significant degree of channeling (χFe ≈ 46 %) was also observed in the Fe spectrum.

Figure 9 shows bright-field TEM cross-section images (taken at [110]Si zone axis) of the asimplanted sample. One clearly observes (Fig. 9a) the amorphous Si layer of ~ 90 nm produced by 40 keV Fe+ implantation and the crystal-amorphous interface. Figure 9b shows in detail the c-a interface, where one can note a large amount of defects called "end-of-range defects" generated by the implantation process. This defective intermediate zone between

out-of-plane and in-plane was tailored by X-ray Bragg-Surface Diffraction technique.

 **ion implantation in Si(100) and recrystallization process**

**3.1 Fe<sup>+</sup>**

interface).

channel-cut monochromator.

**3.1.1 RBS and TEM – Results and discussion**

the two distinct regions (crystalline and amorphous) is mainly composed of dangling bonds, and this particular kind of structural defect is responsible for the IBIEC process. Upon IBIEC conditions (temperature + irradiation) there is a dynamic rearrangement of these dangling bonds with annihilation in pairs which promotes the layer-by-layer planar recrystallization toward the surface.

Fig. 8. 1 MeV He+ RBS spectra in channeling along the [100] direction from a 90 nm a-Si layer containing Fe before (as-implanted sample) and after (recrystallized sample) irradiation with 600 keV Si+ to a dose of 6x1016 ions/cm2 at 350 °C.

Fig. 9. Bright-field TEM cross-sectional images of the as-implanted sample. **(a)** Overview of the 90 nm thick a-Si layer containing Fe. **(b)** High-resolution of the selected region in **(a)** showing in detail the c-a interface region.

Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications 363

Selected-area electron diffraction (SAED) pattern obtained from a plan-view specimen is presented in figure 10e. In addition to strong Bragg reflections of Si, extra spots due to the nanoparticles are apparent in the diffraction pattern. The extra spots show symmetric net patterns suggesting that there is a certain orientation relationship between the nanoparticles and the substrate lattice. From the analysis of the extra spots symmetry and lattice spacing, the diffraction pattern of Figure 10e can be explained as the overlap of cubic -FeSi2 and tetragonal α-FeSi2 phases. The reflections are consistent with α<sup>1</sup> [112] , α<sup>2</sup> [11 2] (straight lines) and [100]γ (dotted line) net patterns (Behar et al., 1996; Vouroutzis et al., 2008). As the SAED measurements have identified two crystalline phases, the near-surface precipitated

Bragg-Surface Diffraction (BSD) (Chang, 2004) is a special diffraction case of the X-ray multiple diffraction (XRMD) technique which has become a very useful and high resolution probe to study in-plane effects in single crystals in general, and also, with several interesting contributions to semiconductor epitaxial systems (dos Santos et al., 2009; Morelhão & Cardoso, 1993; Morelhão et al., 1998; Orloski et al., 2005; Lang et al., 2010b). For a more complete understanding of the experimental results that will follow in this chapter, we

The XRMD phenomenon is systematically generated by aligning the primary planes of a single crystal - generally parallel to the sample surface, to diffract the incident beam and, by rotating it around the normal to the primary (*h*p, *k*p, *l*p) planes while the diffracted beam is monitored by a detector. Under rotation (-axis), several other secondary (*h*s, *k*s, *l*s) planes which are inclined with respect to the surface can enter into diffraction condition simultaneously with the primary ones. A closer observation of the diffraction geometry shows that other diffraction planes, the so-called coupling (*h*c, *k*c, *l*c), also interact with the secondary diffracted beams to re-scatter them towards the detector. The obtained XRMD pattern, called Renninger scanning (RS) (Renninger, 1937), shows a series of peaks distributed according to the symmetry of the chosen primary vector and also to the symmetry plane established by rotation of the several reciprocal space secondary points when entering and leaving the Ewald sphere. Therefore, one can clearly observe in a RS, these two types of symmetry mirrors whose position and intensity distributions are essential for most of the applications of this technique. When a peak in the RS represents an interaction of the incident, the primary and one secondary beam, it shows up as a threebeam peak (or three-beam case). However, one can have two or three secondary beams simultaneously interacting to provide four or five-beam cases (or even cases for *n* > 5 interacting beams) being these secondary beams either Bragg (reflected) or Laue

A special three-beam XRMD case, called Bragg-Surface Diffraction (BSD), appears under adequate conditions, that is, when the secondary diffracted beam propagates along the crystal surface under an extreme asymmetric geometry. A schematic diagram of the multiple scattering for the BSD case occurring inside the crystal, can be seen in Figure 11, where **H**ij are the reciprocal lattice vectors corresponding to the primary planes (**H**01), secondary planes (**H**02) and coupling plane (**H**21). These BSD beams carry information on the

layer should contain α-FeSi2 nanoparticles.

briefly discuss the physical aspects of the XRMD technique.

**3.2 X-ray multiple diffraction**

(transmitted) cases.

The complete recrystallization of the amorphous Si layer was also confirmed by TEM analyses. It should be noted that for a conventional thermal annealing at the same temperature (350 °C) and time (irradiation time ~ 160 minutes) would have produced a regrowth of only 7.7x10-3 nm which is a negligible amount. Therefore, ion-beam irradiation strongly enhances the kinetics of recrystallization. Figure 10 exhibits representative TEM micrographs of the recrystallized sample. The cross-section image, such as Fig. 10a taken along the [110]Si pole and slightly tilted on the zone axis, revealed the efficient a-Si regrowth and an impurity redistribution - nanoparticles formation after the IBIEC process. Iron was completely swept by the moving c-a interface and retained within the precipitated narrowing layer. Despite the amount of Fe present in the recrystallized region, the quality of the recovered crystal appears to be very good, as demonstrated by the TEM image. Three regions regarding the nanoparticles distribution are observed: a thin region a few nanometers thick which is closer to the surface (R1); a Si region about 5 nm with a small concentration of nanoparticles (R2); and a layer (≈ 40 nm wide) with a higher concentration of nanoparticles (R3).

High-resolution cross-sectional images (HRTEM) of the R1 and R3 regions are shown in the insets 10b, 10c and 10d. In inset 10b, it is possible to identify small irregular shaped nanoparticles at Si subsurface R1. In the deeper layers (R2 and R3), two morphological variants of the metastable -FeSi2 phase were observed and recognized: spherical-like nanoparticles epitaxially formed in the substrate with a fully aligned orientation regarding the Si matrix (Fig. 10c) and plate-like nanoparticles rotated with respect to the Si matrix (Fig. 10d) as previously reported (Lin et al., 1994). The spherical-like nanoparticles form coherent interfaces with the Si matrix, while the plate-like ones are elongated along Si 112 directions.

Fig. 10. TEM images of the recrystallized sample. **(a)** [110]Si cross-sectional revealing three nanoparticles regions in depth. HRTEM showing: **(b)** irregular shaped α-FeSi2 nanoparticles at Si subsurface, **(c)** γ-FeSi2 spherical-like and **(d)** γ-FeSi2 plate-like nanoparticles in a deeper region. **(e)** [001]Si SAED pattern indicating the concomitant presence of both α- and γ-FeSi2 phases.

Selected-area electron diffraction (SAED) pattern obtained from a plan-view specimen is presented in figure 10e. In addition to strong Bragg reflections of Si, extra spots due to the nanoparticles are apparent in the diffraction pattern. The extra spots show symmetric net patterns suggesting that there is a certain orientation relationship between the nanoparticles and the substrate lattice. From the analysis of the extra spots symmetry and lattice spacing, the diffraction pattern of Figure 10e can be explained as the overlap of cubic -FeSi2 and tetragonal α-FeSi2 phases. The reflections are consistent with α<sup>1</sup> [112] , α<sup>2</sup> [11 2] (straight lines) and [100]γ (dotted line) net patterns (Behar et al., 1996; Vouroutzis et al., 2008). As the SAED measurements have identified two crystalline phases, the near-surface precipitated layer should contain α-FeSi2 nanoparticles.

#### **3.2 X-ray multiple diffraction**

362 Recrystallization

The complete recrystallization of the amorphous Si layer was also confirmed by TEM analyses. It should be noted that for a conventional thermal annealing at the same temperature (350 °C) and time (irradiation time ~ 160 minutes) would have produced a regrowth of only 7.7x10-3 nm which is a negligible amount. Therefore, ion-beam irradiation strongly enhances the kinetics of recrystallization. Figure 10 exhibits representative TEM micrographs of the recrystallized sample. The cross-section image, such as Fig. 10a taken along the [110]Si pole and slightly tilted on the zone axis, revealed the efficient a-Si regrowth and an impurity redistribution - nanoparticles formation after the IBIEC process. Iron was completely swept by the moving c-a interface and retained within the precipitated narrowing layer. Despite the amount of Fe present in the recrystallized region, the quality of the recovered crystal appears to be very good, as demonstrated by the TEM image. Three regions regarding the nanoparticles distribution are observed: a thin region a few nanometers thick which is closer to the surface (R1); a Si region about 5 nm with a small concentration of nanoparticles (R2); and a layer (≈ 40 nm wide) with a higher concentration

High-resolution cross-sectional images (HRTEM) of the R1 and R3 regions are shown in the insets 10b, 10c and 10d. In inset 10b, it is possible to identify small irregular shaped nanoparticles at Si subsurface R1. In the deeper layers (R2 and R3), two morphological variants of the metastable -FeSi2 phase were observed and recognized: spherical-like nanoparticles epitaxially formed in the substrate with a fully aligned orientation regarding the Si matrix (Fig. 10c) and plate-like nanoparticles rotated with respect to the Si matrix (Fig. 10d) as previously reported (Lin et al., 1994). The spherical-like nanoparticles form coherent interfaces with the Si matrix, while the plate-like ones are elongated along Si 112

Fig. 10. TEM images of the recrystallized sample. **(a)** [110]Si cross-sectional revealing three nanoparticles regions in depth. HRTEM showing: **(b)** irregular shaped α-FeSi2 nanoparticles at Si subsurface, **(c)** γ-FeSi2 spherical-like and **(d)** γ-FeSi2 plate-like nanoparticles in a deeper region. **(e)** [001]Si SAED pattern indicating the concomitant presence of both α- and γ-FeSi2

of nanoparticles (R3).

directions.

phases.

Bragg-Surface Diffraction (BSD) (Chang, 2004) is a special diffraction case of the X-ray multiple diffraction (XRMD) technique which has become a very useful and high resolution probe to study in-plane effects in single crystals in general, and also, with several interesting contributions to semiconductor epitaxial systems (dos Santos et al., 2009; Morelhão & Cardoso, 1993; Morelhão et al., 1998; Orloski et al., 2005; Lang et al., 2010b). For a more complete understanding of the experimental results that will follow in this chapter, we briefly discuss the physical aspects of the XRMD technique.

The XRMD phenomenon is systematically generated by aligning the primary planes of a single crystal - generally parallel to the sample surface, to diffract the incident beam and, by rotating it around the normal to the primary (*h*p, *k*p, *l*p) planes while the diffracted beam is monitored by a detector. Under rotation (-axis), several other secondary (*h*s, *k*s, *l*s) planes which are inclined with respect to the surface can enter into diffraction condition simultaneously with the primary ones. A closer observation of the diffraction geometry shows that other diffraction planes, the so-called coupling (*h*c, *k*c, *l*c), also interact with the secondary diffracted beams to re-scatter them towards the detector. The obtained XRMD pattern, called Renninger scanning (RS) (Renninger, 1937), shows a series of peaks distributed according to the symmetry of the chosen primary vector and also to the symmetry plane established by rotation of the several reciprocal space secondary points when entering and leaving the Ewald sphere. Therefore, one can clearly observe in a RS, these two types of symmetry mirrors whose position and intensity distributions are essential for most of the applications of this technique. When a peak in the RS represents an interaction of the incident, the primary and one secondary beam, it shows up as a threebeam peak (or three-beam case). However, one can have two or three secondary beams simultaneously interacting to provide four or five-beam cases (or even cases for *n* > 5 interacting beams) being these secondary beams either Bragg (reflected) or Laue (transmitted) cases.

A special three-beam XRMD case, called Bragg-Surface Diffraction (BSD), appears under adequate conditions, that is, when the secondary diffracted beam propagates along the crystal surface under an extreme asymmetric geometry. A schematic diagram of the multiple scattering for the BSD case occurring inside the crystal, can be seen in Figure 11, where **H**ij are the reciprocal lattice vectors corresponding to the primary planes (**H**01), secondary planes (**H**02) and coupling plane (**H**21). These BSD beams carry information on the

Ion-Beam-Induced Epitaxial Recrystallization Method and Its Recent Applications 365

Fig. 12. High resolution rocking curves of the recrystallized sample for (004) reflection at = 0º **(a)** and = 90º **(b)** and (002) reflection at two BSD peaks: 6.04º **(c)** and 83.96º **(d)**. Si

As (002) is a forbidden reflection of the Si space group, no primary intensity can be observed out of the BSD secondary peaks. The rocking obtained at 6.04° shows three different contributions: the stronger peak due to the Si matrix; the broad peak to the right (higher angles), due to the R2 and R3 convoluted peaks; and a distinct peak to the left (lower angles), probably associated only with the R3 region. It should be noticed that the rocking at 83.96° exhibits a meaningful profile difference, that is, the peak to the left (lower angles) does not appear as discriminated as in the previous measurement ( 6.04°), it means, a noticeable confirmation of the anisotropic behavior. This anisotropy, observed between the [ 110 ] and [110] in-plane directions, could be associated with the plate-like nanoparticles (Fig. 10d) since the shape and orientation of these ordered nanoparticles should introduce

Figure 13 shows the measured : mappings for a Si matrix (pristine) and a recrystallized sample for comparison purposes to provide a better visualization and characterization of the detected anisotropy. The exact BSD reflection is tailored in both and -directions for each of the two above mentioned BSD secondary reflections: ( 111 ) Si pristine in Fig. 13a and recrystallized sample in 13b and, for (111) Si pristine in 13c and recrystallized one in 13d. In fact, these mappings give a more complete 3D view of the BSD reflection condition which complements the 2D analysis obtained from the HRRC in Figs. 12c and 12d. Furthermore, the mappings allow for the lattice parameters and 2D-strain determination of the distorted regions (R2 and R3). The mappings obtained for the Si matrix (pristine) along the two inplane perpendicular directions as depicted in 13a and 13c, exhibit only the BSD matrix peak at ω = 16.578° and, as expected, no difference is observed. In turn, the recrystallization process induces FeSi2 nanoparticles nucleation within the implanted matrix and then, a huge broadening as well as a striking difference is clearly observed in Figs. 13b and 13d mappings. Besides the BSD matrix peak in Fig. 13b, two other peaks are also detected: one

pristine is also added for comparison purposes.

different strains in both perpendicular directions.

sample surface which are useful for studying surface impurity incorporation effects (Lai et al., 2005), or even on the interface (layer/substrate), if such an interface is present, as in the semiconductor epitaxial structures. The technique has provided significant information when successfully applied as a special high-resolution 3D probe for the study of epitaxial layered heterostructures. Here, the layer and substrate lattices may be separately studied just by selecting one adequate layer or substrate peak. The technique is also a method capable of measuring strain fields at the interfaces in epilayer/substrate systems with depth penetration of 2 Å resolution and with enough sensitivity to detect lattices distortions in the range of 72 Å around the epilayer/substrate interface (Sun et al., 2006).

Fig. 11. BSD scheme using its consecutive scattering model with **H**01 (primary), **H**<sup>02</sup> (secondary) and **H**21 (coupling) vectors. The coupling planes re-scatter the secondary beam towards the primary direction.

Besides the standard RS, another XRMD scanning methodology can give information on the crystalline quality obtained from an analysis of the ω: mapping scans (Morelhão & Cardoso, 1996). By using this method, the rotation is performed for a range of angles each targeting an exact angular position of the multiple-beam Bragg condition. This approach results in a three dimensional plot of the primary intensity versus and in a coupled way from which, through analysis of the iso-intensity contours of such plots, one can obtain information on the lattice coherence along the beam path and hence, on the crystalline quality. It has been shown that when the FWHM (full width at half maximum) of the peak in the scan is larger in comparison to the one in the scan, there is almost no loss of coherence, i.e. confirming that the crystal is perfect or nearly perfect.

#### **3.2.1 BSD reflections – Results and discussion**

Measurements of the (004) symmetrical high resolution rocking curves (HRRC) are shown in figure 12 at two perpendicular orientations on the sample surface: 0° (Fig. 12a) and 90° (Fig. 12b). Both patterns present practically the same result with two distinct peaks corresponding to R2 and R3 regions, clearly seen in each pattern, with smaller perpendicular lattice parameters in comparison to the matrix peak (stronger). Also, (002) HRRC were measured at ( 111 ) and (111) BSD reflections then, at two azimuth angles on the recrystallized sample ( 6.04° and 83.96°). The results are shown in figures 12c and 12d.

sample surface which are useful for studying surface impurity incorporation effects (Lai et al., 2005), or even on the interface (layer/substrate), if such an interface is present, as in the semiconductor epitaxial structures. The technique has provided significant information when successfully applied as a special high-resolution 3D probe for the study of epitaxial layered heterostructures. Here, the layer and substrate lattices may be separately studied just by selecting one adequate layer or substrate peak. The technique is also a method capable of measuring strain fields at the interfaces in epilayer/substrate systems with depth penetration of 2 Å resolution and with enough sensitivity to detect lattices distortions in the

range of 72 Å around the epilayer/substrate interface (Sun et al., 2006).

Fig. 11. BSD scheme using its consecutive scattering model with **H**01 (primary), **H**<sup>02</sup>

of coherence, i.e. confirming that the crystal is perfect or nearly perfect.

**3.2.1 BSD reflections – Results and discussion**

towards the primary direction.

and 12d.

(secondary) and **H**21 (coupling) vectors. The coupling planes re-scatter the secondary beam

Besides the standard RS, another XRMD scanning methodology can give information on the crystalline quality obtained from an analysis of the ω: mapping scans (Morelhão & Cardoso, 1996). By using this method, the rotation is performed for a range of angles each targeting an exact angular position of the multiple-beam Bragg condition. This approach results in a three dimensional plot of the primary intensity versus and in a coupled way from which, through analysis of the iso-intensity contours of such plots, one can obtain information on the lattice coherence along the beam path and hence, on the crystalline quality. It has been shown that when the FWHM (full width at half maximum) of the peak in the scan is larger in comparison to the one in the scan, there is almost no loss