*1.4.2 Effect of the growth temperature on the value of tensile strain*

After the low-temperature growth step, we have deposited another Ge film at higher substrate temperatures. In the following, we will consider the effect of the two-step growth process on the tensile strain in the Ge films. All samples have a total thickness of ~ 300 nm, which consists of a ~50 nm thick Ge film deposited at 300°C followed by a 250 nm thick Ge layer grown at various temperatures: 400, 500, 600, 650, 700, 750 and 770°C. As the tensile strain in Ge layers is induced by the difference of the thermal expansion coefficients between Ge and Si, it is natural to expect that the level of tensile strain should increase with increasing the growth temperature. **Figure 7** displays some representative θ–2θ XRD scans taken around the Ge(004) reflection. For comparison, we report in dotted lines a XRD scan of a sample grown at 300°C of the same thickness.

As the growth temperature increases, the Ge(004) peak linearly shifts to higher angles, reaches a saturation value at 700°C and finally remains almost constant for further increasing the temperature to 770°C. The in-plane tensile strain observed in the temperature range of 700–770°C is 0.24%. The fact that the (004) peak of the 300°C grown sample is located at 2θ ~ 66°, a value close to that measured on a Ge substrate, indicates that the corresponding Ge layer is almost fully relaxed. By using a high-resolution XRD, we estimate the rate of strain relaxation is about 95–96%. As can be seen in **Figure 7**, the 2θ angle value of the (004) peak is found to linearly increase when the growth temperature increases from 300 to 700°C then remains almost constant for further increasing the temperature up to 770°C. We note that to obtain the value of the in-plane tensile strain ε//, we first determine the out-of plane strain ε⊥ from the θ–2θ XRD curves and then deduce the value of ε// using the following relationship: ε///(ε// + ε⊥) = c11/(c11 + 2c12) with c11 = 12.85 × 1010 Pa and c12 = 4.83 × 1010 Pa for pure Ge [28]. The highest value of the in-plane tensile strain ε// obtained in the growth temperature range of 700–770°C is 0.24%, which is in agreement with previous results reported using CVD in which the highest tensile strain was in the range of 0.22–0.25% [24–26, 29–33].

As we have already mentioned previous part, one of the unique properties of Ge is the very small energy difference between its direct and indirect band gap. Thus, Si/Ge interdiffusion represents an obstacle to overcome in order to develop Ge-based optoelectronic devices. Si/Ge interdiffusion has been shown to greatly affect the optical properties of Si/Ge heterostructures [34] and degrade the performance of metal-oxide semiconductor field-effect transistors (MOSFET) by reducing strain and carrier confinement and increasing alloy scattering [35]. To prevent interdiffusion or out-diffusion of an element (occurs when the sample was annealed at high temperature to enhance the tensile strain value in the Ge film), it is common to use a diffusion barrier and materials must be not only nonreactive but also are able to strongly adhere to adjacent materials. In electronic or memory devices, multilayers of metals, WN2, RuTiN, or RuTiO, are usually used to prevent out diffusion of dopants (B and P) or oxidation of devices [36–38]. Such materials are, however, difficult insert in a heterostructure where epitaxial growth is needed. It is now generally accepted that Si and Ge atoms interdiffuse via both vacancy and interstitial-related mechanisms [39]. To prevent Si/Ge interdiffusion, in particular to suppress upward diffusion of Si to the deposited Ge layer during growth or subsequent annealing, we have experimented an approach to saturate vacancies and

#### **Figure 7.**

*Evolution of θ–2θ XRD scans around the Ge(004) reflection with the growth temperature. The doted XRD scan corresponding to a sample grown at 300°C is shown for comparison.*

**79**

**Figure 8.**

*New Material for Si-Based Light Source Application for CMOS Technology*

and can easily diffuse via interstitial mechanism [40–42].

interstitial sites in the Ge layer near the interface region. We have chosen carbon for its small atomic radius (twice as small as Ge), carbon atoms are thus highly mobile

To further confirm the effect of Ge/Si diffusion, we have used high-resolution

**Figure 9(a)** shows a TEM image taken near the interface region of a sample in which we have inserted three separate carbon layers during the first step of Ge growth. Starting from a clean and (2 × 1) reconstructed Si surface, we first grow a 30 nm thick Ge buffer layer at 300°C and then about 4 monolayers (ML) of carbon at the same temperature. Since epitaxial growth of the upper Ge layers should be conserved, the amount of deposited carbon is crucial. The latter can be determined using the change of RHEED patterns. Upon carbon deposition at 300°C, the (2 × 1) RHEED pattern characteristic of a clean Ge(001) surface remains almost unchanged up to carbon coverage of 6 MLs, beyond which a faint pattern appears. Since it is crucial to insure epitaxial growth of the upper Ge layer, a carbon amount of 4 MLs is then chosen (the corresponding carbon thickness is ~0.3 nm). After this step, a Ge layer is deposited on top of carbon, producing therefore C/Ge stacked layers in order to increase the efficiency of carbon-induced diffusion barriers. We have experimented three multilayers of C (0.3 nm)/Ge (18 nm). **Figure 9(b)** shows

*Values of inter-reticular distances of Si and Ge, measured along the (004) and (220) planes of a 300 nm thick* 

*Ge layer deposited on Si at 730°C, followed by 10 cyclic anneals at 780/900°C for 20 min.*

HR-XRD reciprocal space mapping (RSM) equipped with a rotating anode to determine the Si composition and the strain level in Ge layers. 004, 224, and −2−24 RSMs were collected along the four <110> azimuthal directions. Combining the (224) map with that along (004) direction, one can measure both parallel and perpendicular lattice parameters of the Ge layer. The data obtained in reciprocal space are converted in direct space. We show in **Figure 8** the values of interreticular distances of the Si substrate and the Ge layer, measured along the (004) and (220) planes, of the sample annealed by 10 cycles at 780/900°C for 20 min. Note that for the as-grown sample, we obtain a// = 5.669 Å and a⊥ = 5.649 Å, which correspond the value of pure Ge (aGe = 5.657 Å). For the annealed sample, we obtain a// = 5.656 Å and a⊥ = 5.625 Å. Thus, the volume occupied by this unit cell corresponds to a SixGe1−x alloy with a Si average concentration x = 5% (the corresponding

*DOI: http://dx.doi.org/10.5772/intechopen.84994*

average lattice parameter is about 5.646 Å).

#### *New Material for Si-Based Light Source Application for CMOS Technology DOI: http://dx.doi.org/10.5772/intechopen.84994*

*Silicon Materials*

As the growth temperature increases, the Ge(004) peak linearly shifts to higher angles, reaches a saturation value at 700°C and finally remains almost constant for further increasing the temperature to 770°C. The in-plane tensile strain observed in the temperature range of 700–770°C is 0.24%. The fact that the (004) peak of the 300°C grown sample is located at 2θ ~ 66°, a value close to that measured on a Ge substrate, indicates that the corresponding Ge layer is almost fully relaxed. By using a high-resolution XRD, we estimate the rate of strain relaxation is about 95–96%. As can be seen in **Figure 7**, the 2θ angle value of the (004) peak is found to linearly increase when the growth temperature increases from 300 to 700°C then remains almost constant for further increasing the temperature up to 770°C. We note that to obtain the value of the in-plane tensile strain ε//, we first determine the out-of plane strain ε⊥ from the θ–2θ XRD curves and then deduce the value of ε// using the following relationship: ε///(ε// + ε⊥) = c11/(c11 + 2c12) with c11 = 12.85 × 1010 Pa and c12 = 4.83 × 1010 Pa for pure Ge [28]. The highest value of the in-plane tensile strain ε// obtained in the growth temperature range of 700–770°C is 0.24%, which is in agreement with previous results reported using CVD in which the highest tensile

As we have already mentioned previous part, one of the unique properties of Ge is the very small energy difference between its direct and indirect band gap. Thus, Si/Ge interdiffusion represents an obstacle to overcome in order to develop Ge-based optoelectronic devices. Si/Ge interdiffusion has been shown to greatly affect the optical properties of Si/Ge heterostructures [34] and degrade the performance of metal-oxide semiconductor field-effect transistors (MOSFET) by reducing strain and carrier confinement and increasing alloy scattering [35]. To prevent interdiffusion or out-diffusion of an element (occurs when the sample was annealed at high temperature to enhance the tensile strain value in the Ge film), it is common to use a diffusion barrier and materials must be not only nonreactive but also are able to strongly adhere to adjacent materials. In electronic or memory devices, multilayers of metals, WN2, RuTiN, or RuTiO, are usually used to prevent out diffusion of dopants (B and P) or oxidation of devices [36–38]. Such materials are, however, difficult insert in a heterostructure where epitaxial growth is needed. It is now generally accepted that Si and Ge atoms interdiffuse via both vacancy and interstitial-related mechanisms [39]. To prevent Si/Ge interdiffusion, in particular to suppress upward diffusion of Si to the deposited Ge layer during growth or subsequent annealing, we have experimented an approach to saturate vacancies and

*Evolution of θ–2θ XRD scans around the Ge(004) reflection with the growth temperature. The doted XRD scan* 

*corresponding to a sample grown at 300°C is shown for comparison.*

strain was in the range of 0.22–0.25% [24–26, 29–33].

**78**

**Figure 7.**

interstitial sites in the Ge layer near the interface region. We have chosen carbon for its small atomic radius (twice as small as Ge), carbon atoms are thus highly mobile and can easily diffuse via interstitial mechanism [40–42].

To further confirm the effect of Ge/Si diffusion, we have used high-resolution HR-XRD reciprocal space mapping (RSM) equipped with a rotating anode to determine the Si composition and the strain level in Ge layers. 004, 224, and −2−24 RSMs were collected along the four <110> azimuthal directions. Combining the (224) map with that along (004) direction, one can measure both parallel and perpendicular lattice parameters of the Ge layer. The data obtained in reciprocal space are converted in direct space. We show in **Figure 8** the values of interreticular distances of the Si substrate and the Ge layer, measured along the (004) and (220) planes, of the sample annealed by 10 cycles at 780/900°C for 20 min. Note that for the as-grown sample, we obtain a// = 5.669 Å and a⊥ = 5.649 Å, which correspond the value of pure Ge (aGe = 5.657 Å). For the annealed sample, we obtain a// = 5.656 Å and a⊥ = 5.625 Å. Thus, the volume occupied by this unit cell corresponds to a SixGe1−x alloy with a Si average concentration x = 5% (the corresponding average lattice parameter is about 5.646 Å).

**Figure 9(a)** shows a TEM image taken near the interface region of a sample in which we have inserted three separate carbon layers during the first step of Ge growth. Starting from a clean and (2 × 1) reconstructed Si surface, we first grow a 30 nm thick Ge buffer layer at 300°C and then about 4 monolayers (ML) of carbon at the same temperature. Since epitaxial growth of the upper Ge layers should be conserved, the amount of deposited carbon is crucial. The latter can be determined using the change of RHEED patterns. Upon carbon deposition at 300°C, the (2 × 1) RHEED pattern characteristic of a clean Ge(001) surface remains almost unchanged up to carbon coverage of 6 MLs, beyond which a faint pattern appears. Since it is crucial to insure epitaxial growth of the upper Ge layer, a carbon amount of 4 MLs is then chosen (the corresponding carbon thickness is ~0.3 nm). After this step, a Ge layer is deposited on top of carbon, producing therefore C/Ge stacked layers in order to increase the efficiency of carbon-induced diffusion barriers. We have experimented three multilayers of C (0.3 nm)/Ge (18 nm). **Figure 9(b)** shows

#### **Figure 8.**

*Values of inter-reticular distances of Si and Ge, measured along the (004) and (220) planes of a 300 nm thick Ge layer deposited on Si at 730°C, followed by 10 cyclic anneals at 780/900°C for 20 min.*

**Figure 9.**

*(a) TEM image taken near the interface region illustrating the growth of three multilayers of C (0.3 nm)/Ge (18 nm); (b) atomically resolved TEM image taken in the vicinity of the carbon layer.*

an atomically resolved TEM image taken in the vicinity of the carbon layer. Clearly, the underneath Ge layers and also the upper Ge layers are perfectly epitaxial, without any detectable defects. The image also reveals that carbon atoms are distributed over a distance of ~2 nm, probably to occupy the interstitial sites of the Ge lattice or due to the strain generation arising from carbon insertion.

To verify the efficiency of C/Ge multilayers to prevent Si/Ge interdiffusion, we have finally grown a 300 nm thick Ge layer following the two-step growth as described previously. The high temperature step was carried out at 730°C. After growth, the sample was cyclically annealed by 10 cycles from 780 to 900°C and the annealing time was 20 min. In **Figure 10**, we display SIMS measurements of two samples: one without carbon deposition (red curve) and the other containing three C/Ge multilayers. As can be seen, the as-grown sample (black) shows a relatively sharp interface, no long-range upward diffusion of Si from the substrate was observed and the deposited Ge layer remains pure. On the other hand, the Si profile of the annealed sample, indicated in blue, reveals a pronounced Si upward diffusion. The Si content within the Ge film is highly heterogeneous, it continuously decreases from the interface to the film surface. The average Si content in the Ge film is estimated to be about 5%.

#### **1.5 Phosphorus doping using a GaP decomposition source**

Heavy n-doping of Ge films is essential to achieve efficient light emission from the direct gap transition. However, heavy n-type doping in Ge is a challenge due to a low solubility and a fast diffusivity of dopants. Three elements: phosphorus, arsenic and antimony, can be used as n-type doping in group-IV semiconductors. We report in **Table 2** the solubility of these three elements in Ge [43].

It can be seen that phosphorous has the highest solubility and at a temperature not too high. In UHV-CVD, the highest phosphorus concentration of 1.2 × 1019 cm−<sup>3</sup>

**81**

*New Material for Si-Based Light Source Application for CMOS Technology*

was reported [44]. In CVD, the PH3 is commonly used as a gas precursor and

*SIMS profiles measuring the Si content of the as-grown sample (black) and after 10 cyclic anneals at 780/900°C for 20 min (blue). The red curve corresponds to a sample in which three short C/Ge multilayers were deposited* 

by combining delta doping with diffusion barriers to reduce the P out-diffusion [45, 46]. It is worth noting that while dopant implantation is commonly used in the CMOS process, the activated P concentration is limited at about 1 × 1019 cm−<sup>3</sup> even with the use of diffusion barriers [47], and a higher P concentration seriously deteriorates the material quality. Since tetrahedral white phosphorus (P4 molecules) is not stable and highly volatile, we have used a specific doping cell based on the decomposition of GaP to produce di-phosphorus (P2 molecules) [46–48], which has a sticking coefficient of about 10 times larger than that of tetrahedral white phosphorus. Thus, the phosphorus concentration in the Ge lattice could be increased and we can expect to enhance the efficiency of radiative recombination in Ge.

A photograph of a GaP decomposition cell is shown in **Figure 11**, which is heated by a Ta wire filament supported by PBN rings, similar to standard effusion cells. Compared to a valved phosphorus thermal cracker, the GaP cell is easy to set up, it operates as any effusion cell and is compatible with MBE environments. Using a Ga-trapping cap system, the GaP source can provide high-purity P2 beam by decomposition of GaP. Its operation is based on the sublimation of phosphorus from GaP at an intermediate temperature range in which Ga has a very low vapor pressure. By placing a cap with small apertures on top of the cell (**Figure 11(b)**),

600–700°C [12]. A phosphorus concentration up to 4.5 × 1019 cm−<sup>3</sup>

was obtained at a temperatures of

can be achieved

an activated P concentration of 2 × 1019 cm−<sup>3</sup>

*Solubility of different n-type dopants in Ge [43].*

*1.5.1 Specific GaP decomposition source*

*DOI: http://dx.doi.org/10.5772/intechopen.84994*

**Figure 10.**

**Table 2.**

*near the interface region.*

*New Material for Si-Based Light Source Application for CMOS Technology DOI: http://dx.doi.org/10.5772/intechopen.84994*

#### **Figure 10.**

*Silicon Materials*

**Figure 9.**

an atomically resolved TEM image taken in the vicinity of the carbon layer. Clearly, the underneath Ge layers and also the upper Ge layers are perfectly epitaxial, without any detectable defects. The image also reveals that carbon atoms are distributed over a distance of ~2 nm, probably to occupy the interstitial sites of the Ge lattice or

*(a) TEM image taken near the interface region illustrating the growth of three multilayers of C (0.3 nm)/Ge* 

To verify the efficiency of C/Ge multilayers to prevent Si/Ge interdiffusion, we have finally grown a 300 nm thick Ge layer following the two-step growth as described previously. The high temperature step was carried out at 730°C. After growth, the sample was cyclically annealed by 10 cycles from 780 to 900°C and the annealing time was 20 min. In **Figure 10**, we display SIMS measurements of two samples: one without carbon deposition (red curve) and the other containing three C/Ge multilayers. As can be seen, the as-grown sample (black) shows a relatively sharp interface, no long-range upward diffusion of Si from the substrate was observed and the deposited Ge layer remains pure. On the other hand, the Si profile of the annealed sample, indicated in blue, reveals a pronounced Si upward diffusion. The Si content within the Ge film is highly heterogeneous, it continuously decreases from the interface to the film surface. The average Si content in the Ge

Heavy n-doping of Ge films is essential to achieve efficient light emission from the direct gap transition. However, heavy n-type doping in Ge is a challenge due to a low solubility and a fast diffusivity of dopants. Three elements: phosphorus, arsenic and antimony, can be used as n-type doping in group-IV semiconductors. We report

It can be seen that phosphorous has the highest solubility and at a temperature not too high. In UHV-CVD, the highest phosphorus concentration of 1.2 × 1019 cm−<sup>3</sup>

due to the strain generation arising from carbon insertion.

*(18 nm); (b) atomically resolved TEM image taken in the vicinity of the carbon layer.*

**1.5 Phosphorus doping using a GaP decomposition source**

in **Table 2** the solubility of these three elements in Ge [43].

film is estimated to be about 5%.

**80**

*SIMS profiles measuring the Si content of the as-grown sample (black) and after 10 cyclic anneals at 780/900°C for 20 min (blue). The red curve corresponds to a sample in which three short C/Ge multilayers were deposited near the interface region.*


#### **Table 2.**

*Solubility of different n-type dopants in Ge [43].*

was reported [44]. In CVD, the PH3 is commonly used as a gas precursor and an activated P concentration of 2 × 1019 cm−<sup>3</sup> was obtained at a temperatures of 600–700°C [12]. A phosphorus concentration up to 4.5 × 1019 cm−<sup>3</sup> can be achieved by combining delta doping with diffusion barriers to reduce the P out-diffusion [45, 46]. It is worth noting that while dopant implantation is commonly used in the CMOS process, the activated P concentration is limited at about 1 × 1019 cm−<sup>3</sup> even with the use of diffusion barriers [47], and a higher P concentration seriously deteriorates the material quality. Since tetrahedral white phosphorus (P4 molecules) is not stable and highly volatile, we have used a specific doping cell based on the decomposition of GaP to produce di-phosphorus (P2 molecules) [46–48], which has a sticking coefficient of about 10 times larger than that of tetrahedral white phosphorus. Thus, the phosphorus concentration in the Ge lattice could be increased and we can expect to enhance the efficiency of radiative recombination in Ge.

#### *1.5.1 Specific GaP decomposition source*

A photograph of a GaP decomposition cell is shown in **Figure 11**, which is heated by a Ta wire filament supported by PBN rings, similar to standard effusion cells. Compared to a valved phosphorus thermal cracker, the GaP cell is easy to set up, it operates as any effusion cell and is compatible with MBE environments. Using a Ga-trapping cap system, the GaP source can provide high-purity P2 beam by decomposition of GaP. Its operation is based on the sublimation of phosphorus from GaP at an intermediate temperature range in which Ga has a very low vapor pressure. By placing a cap with small apertures on top of the cell (**Figure 11(b)**),

**Figure 11.** *(a) Photograph of a GaP decomposition cell and (b) schema of the Ga trapping cap [50].*

Ga atoms are trapped by the cap and only the P2 beam can escape. According to the supplier [49], a P2/P4 ratio of about 150:1 can be achieved.
