Ge Growth on Si(001)

Due to the ~4.1% lattice mismatch between Si (5.431 Ǻ) and Ge (5.658 Ǻ), hetero-epitaxial growth of Ge on Si beyond critical thickness generally results in the formation of misfits and associated threading dislocations even islanding .There\are a number of ways to grow high quality single crystalline Ge on Si including using a graded GexSi1-x buffer [8, 9] or overgrowing Ge in a Si/SiO2 template [10]. Here we use a two-step growth approach [11, 12] with cycling growth and anneal to grow high quality Ge on Si substrate.

The low-temperature buffer layer is to prevent the Ge from Stranski-Krastanov (S-K) Fig.4-12 The SiGe growth rate and Ge content using (a) SiH4/GeH4 with H2 and

growth. Base on S-K growth theory, only Ge layer less than its critical thickness (~10¹ nm) can be free from dislocations and induce strain (compressive strain in this case) in Ge. Above critical thickness, the energy of the strain stored in the Ge can be relaxed by the formation of misfit dislocation (a plastic process) or three-dimensional island. To prevent such island growth, the temperature must be low enough to kinetically prohibit adatoms moving to form islands with adequate kinetic energy. Besides, hydrogen on the surface can act as surfactant at low temperature and can hinder the nucleation of the 3-dimensional islands [13]. In practice, the low temperature can also prevent the complete equilibrium condition where dislocation may start to form even before the critical thickness. Once the strain energy is fully relaxed plastically by misfit dislocation, raised Ge growth temperature for growth rate at higher temperature can have on the low-temperature Ge buffer layer, which is in the homo-epitaxial case. The low temperature 320^oC with 375^oC is used for low and high growth temperature for our Ge, respectively.

Cycling thermal annealing at 825^oC after growth is performed to reduce decrease threading dislocations which propagate to the edge of the substrate at annealing. In the ref. 14, 1-2 order of reduction of threading dislocation density can have using thermal annealing at 900^oC. These threading dislocations act as both defect centers (non-radiative recombination centers) which degrade the electrical and optical properties of material. Fig. 4-13 is the 4.1mm epitaxial Ge on Si(001) with the 2.7x10⁶ cm^-2 dislocation density which is observed by etching pits using Schimmel etching. Due to the larger thermal expansion coefficient of Ge than Si, Ge tends to shrink more by

post annealing and induces ~0.16% in-plane tensile strain in Ge film upon cooling (Fig.

4-14).

The room temperature PL is taken for comparison with bulk Ge. The PL peaks at 695 meV and 780 meV are attributed to the indirect band transition and direct band transition, respectively (Fig.4-15). For the radiative recombination in an indirect bandgap material such as Ge, momentum conservation is achieved through the longitudinal acoustic phonon (28 meV) assisted transition [15]. In a direct transition, the momentum conserved is achieved without phonon. The energy difference of conduction band minimum between valley and L valley is 140 meV at room temperature [16].

The spectra of the indirect and direct band gap emission in both bulk Ge and epi-Ge on Si sample were fitted by using the electron-hole plasma recombination model [17] [15]

and the direct band gap recombination model [18], respectively. The band tail of absorption edge model is also taken into consideration in the direct band gap recombination model [19]. The integrated PL intensity ratios of direct to indirect band gap transition for the bulk n-Ge and epi-Ge on Si sample are 0.05 and 2.5, respectively.

For the indirect radiative recombination in bulk Ge, a dominant LA phonon (~28 meV) is involved to satisfy the momentum conservation between the L valleys and the zone center in the valance band.

For the epi-Ge on Si sample, the large defect density in the Ge film may lead to a spread of trap levels in momentum space. These trap levels enhance the non-radiative recombination rate because the more phonons with different momentum can be involved in the momentum conservation. Thus, the relative intensity of the indirect band gap transition is lower in the epi-Ge on Si sample than in the bulk Ge [20].

Threading dislocation density: 2.7 × 10⁶cm^-2

4.1um

Fig.4-13 The high resolution TEM image of 4.1m Ge on Si(001). The inset is the Nomarski micrograph of Ge with Shimmel etching for etching pits observation.

Fig.4-14 The -2 XRD of 4.1 m Ge on Si(001). Note the 0.16% tensile strain in Ge film.

64 65 66 67 68 69 70 71 72

Inten sity (a.u.)

2



(degree)

4.1um Ge/Si X'pert Fitting relaxed Ge

e~0.16%

To observe the carrier gas effect on Ge growth, 375^oC, 40torr with nitrogenand hydrogen are used for comparison. Both samples have post annealing in hydrogen at 825^oC to reduce the threading dislocations. It is found that the Ge growth rate enhanced three times in N2 over in H2 (Fig.4-16 (a)). At 375^oC for Ge growth, it is in the reaction-rate-limited regime, where H desorption more in nitrogen ambient can be the main reason for higher growth rate than in hydrogen. More open sites created in nitrogen ambient also lead to smoother surface than in hydrogen(Fig.4-16 (b)). Using N2 as carrier gas strongly reduces the H2 partial pressure compared to a deposition in hydrogen as carrier gas. The balance between the germane and hydrogen adsorption and the hydrogen desorption is modified depending on the nature of the carrier gas [21].

Fig.4-15 The PL spectra of the bulk Ge and epitaxial Ge on Si at room temperature.

The integrated PL intensity ratios of the direct to indirect band gap transition of the bulk n-Ge and epitaxial Ge on Si sample are 0.05 and 2.5, respectively. The bulk Ge is ~15 times the integrated intensity of PL of the epitaxial Ge sample.

Lower hydrogen partial pressure favors germane adsorption and enhance hydrogen desorption to enhance Ge growth rate.

The surface roughness reduction can be explained in terms of the Ge diffusion barrier as compared to the Ge–H cluster [22]. The effect of attaching H to Ge reduces the with surface RMS roughness of 2nm

(a)

4.5 Summary

To reduce the cost of tool investment for epitaxial growth by CVD, the high growth rate of epitaxial layer is crucial for device development. With the growth model by CVD, the growth rate can be changed by Ge content of SiGe film, precursors, or carrier gas. Due to lower H desorption energy from Ge than from Si, more open sites created leads to higher growth rate when Ge% increasing. When using SiH4/ GeH4 for SiGe films and GeH4 for Ge films growth with N2 instead of H2 as carrier gas. The growth condition having more surface open site in N2 than in H2 is the main reason for growth rate enhancement. In the case using DCS/ GeH4 for SiGe films, the need of H2 for the conversion of adsorbed SiCl2 into Si results to growth rate reduction in N2.

Reference:

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[3] S.M. Gates et al., Appl. Phys. Lett.,58, 2964 (1991).

[4] X. Xiao, C. Liu, J. Sturm, L. Lenchyshyn, and M. Thewalt, Appl. Phys. Lett., vol.

60, p. 1720, (1992).

[5] John C. Bean, PROCEEDINGS OF THE IEEE, VOL. 80, NO 4, APRIL (1992).

[6] P. M. Garonne, J. C. Sturm and P. V. Schwartz, Appl. Phys. Lett., 56 (1990) 1275

[7] W.A.P. CLAASSEN and J. BLOEM, Journal of Crystal Growth 50 N807-815 (1980).

[8] E. A. Fitzgerald, Y.-H. Xie, M. L. Green, D. Brasen, A. R. Kortan, J. Michel, Y.-J.

Mii, and B. E. Weir, Appl. Phys. Lett. 59(7), 811~813 (1991).

[9] S. B. Samavedam and E. A. Fitzgerald, J. Appl. Phys. 81(7), 3108~3116 (1997).

[10] T. A. Langdo, C. W. Leitz, M. T. Currie, E. A. Fitzgerald, A. Lochtefeld, and D. A.

Antoniadis, Appl. Phys. Lett. 75(25), 3700~3702 (2000).

[11] L. Colace, G. Masini, F. Galluzzi, G. Assanto, G. Capellini, L. D. Gaspare, E.

Palange, and F. Evangelisti, Appl. Phys. Lett. 72(24), 3175~3177 (1998).

[12] H.-C. Luan, D. R. Lim, K. K. Lee, K. M. Chen, J. G. Sandland, K. Wada, and L. C.

Kimerling, Appl. Phys. Lett 75(19), 2909~2911 (1999).

[13] D. J. Eaglesham, F. C. Unterwald, and D. C. Jacobson, Phys. Rev. Lett. 70(7), 966~969 (1993).

[14] J. Liu, Ph.D. thesis, Massachusetts Institute of Technolog,(2007).

[15] M. H. Liao, T.-H. Cheng, and C. W. Liu, Appl. Phys. Lett., Vol. 89, 261913, (2006).

[16] S. M. Sze, Physics of Semiconductor Devices, 2nd edition, Wiley, New York, (1981).

[17] C. W. Liu, M. H. Lee, M.-J. Chen, I. C. Lin, and C-F Lin, Appl. Phys. Lett., Vol.

76, pp. 1516-1518, (2000).

[18] P.-S. Kuo, B.-C. Hsu, P.-W. Chen, P. S. Chen, and C. W. Liu, Electrochem.

Solid-State Lett., 7 (10), G201, (2004).

[19] M. El Kurdi, T. Kociniewski, T.-P. Ngo, J. Boulmer, D. Débarre, P. Boucaud, J. F.

Damlencourt, O. Kermarrec, and D. Bensahel, Appl. Phys. Lett., Vol. 94, 191107, (2009).

[20] S.-R. Jan,C.-Y. Chen, C.-H. Lee, S.-T. Chan, K.-L. Peng, C. W. Liu,,Y. Yamamoto, and B. Tillack, Appl. Phys. Lett 98,141105, (2011).

[21] J. Pejnefors, S.-L. Zhang, H.H. Radamson, J.V. Grahn, M.O¨ stling, J. Appl. Phys.

88 (3),1655,(2000).

[22] Ammar Nayfeh, Chi On Chui, and Krishna C. Saraswat, Takao Yonehara, Appl.

Phys. Lett 85,14, (2004).

[23] S. Horch, H. T. Lorensen, S. Helveg, E. Laegsgaard, I. Stensgaard, K. W. Jacobsen, J. K. Norskov, and F. Basenbacher, Nature (London) 398, 134 (1999).

Chapter 5

Solid Phase Doped Ge by Chemical Vapor Deposition

5.1 Introduction

Ge has been widely investigated for high mobility channels of metal–oxide–

semiconductor field effect transistors (MOSFETs). To realize high performance Ge MOSFETs, the source/drain doping is essential for scaled devices. Ion implantation still has the possibility to create defects even after implant anneal. For phosphorous, the point defects such as Ge vacancies are the main root cause of fast diffusion of dopants [1] and low activation fraction [2]. Although the diffusivity of boron is smaller than phosphorous in Ge, stable defects by implantation-induced crystal disorder remain after long time annealing [3]. All of these defects degrade the Ge diode performance to have non-ideal forward current, low on/off ratio, and high reverse saturation current. The required high temperature of 650^oC ~ 750^oC to release dopants from the solvent volume by spin-on doping on bulk Ge, results in the deep junction depth of ~500nm [4]. The carrier gas (N2 vs. H2) effects on the deposition of boron layers and phosphorous layers on Ge show that H2 retards the incorporation of B2H6 and PH3 [5, 6]. The Si p⁺n

junction by boron layers with the low reverse saturation current was reported previously [7]. Using the patterned wafers by oxide, the boron or phosphorous can be deposited on the oxide holes and doped the Ge selectively if the oxide is thick enough. The oxide even with the boron or phosphorous layer on top can be removed by etching process afterwards. For source and drain application, the doping level should be as high as possible to reduce the parasitic resistance. To obtain the pre-determined level, the top Ge layer can be etched away to leave the residual Ge with the desired doping concentration. In this work, ultra-high vacuum CVD (UHV/CVD) is used to grow boron and phosphorous layers at 450^oC as the solid source for p-type and n-type dopants in Ge, respectively. The well-behaved diffusion process at 450^oC leads to abrupt boron and phosphorous profiles in Ge due to the suppression of defect-assistant diffusion.

5.2 Experiments

After HF dip, the (100) Si substrate is immediately loaded into UHV/CVD system with the cold-wall stainless steel chamber. The base pressure and the growth pressure are maintained below ~10^-9 torr and ~10^-3 torr, respectively. This high vacuum ensures the partial pressures of water vapor, oxygen, and hydrocarbons to be in the 10^-11 torr range and reduces the contamination during the deposition process. Pure germane with the gas flow of 10 sccm and the growth temperature of 420^oC is used for epitaxial Ge growth on Si (100) substrate. It has been shown the threading dislocation density in the Ge epitaxial film can be reduced by 1-2 order of magnitude to 2x10⁷ cm^-2 using a

post-growth thermal annealing at 900 ^oC [29]. The threading dislocations act as both defect centers (non-radiactive recombination centers) which degrade the electrical and optical properties of material. The threading dislocations connecting from the surface to the bottom of the film are also high conductance paths which is the major cause for the leakage current in Ge diode devices. In order to reduce the threading dislocation density in our epitaxial Ge, the in situ annealing at 900^oC for 10 minutes is used to reduce threading dislocation density in epitaxial Ge layer. The epitaxial Ge on Si has the doping concentration on the order of 1x10¹⁶ cm^-3.After the growth of epitaxial Ge, 0.1%

diborone (B2H6) in H2 and 100ppm phosphine (PH3) in H2 with the flow rate of 50 sccm are injected into the reaction chamber to grow the p-type and n-type dopant layers, respectively, at 450^oC. To reduce the suppression of the boron and phosphorous adsorption on the hydrogen-terminated Ge surface [5, 6] no additional H2 carrier gas is used in the growth. To avoid desorption and to maintain the low thermal budget, no subsequent annealing is performed. The Ge diodes are formed by mesa structures.

Before the electrodes metallization, the boron layer is etched by hot H3PO4 and phosphorous layer is etched by HNO3 with HF dip and DI water rinse afterward. After etching, Ti was deposited as electrodes. The mesa etching of Ge diode is performed by H2O2. The diodes have the mesa structures with the diameter of 100m. There is no passivation on the sidewalls of the mesa.

5.3 Boron and Phosphorous layer on Ge by Chemical Vapor Deposition

Fig. 5-1 shows the cross-sectional transmission electron microscopy (TEM) images of epitaxial 130nm Ge on Si(001) before and after 900^oC annealing. The epitaxial Ge with the threading dislocation density at 5x10⁸ cm^-2 is demonstrated. After B2H6 dn PH3 injection for two hours at 450^oC, boron and phosphorous layers are deposited on (100) Ge (Fig.5-2). The adsorbed B2H6 forms a ~5nm boron layer on the Ge surface (Fig. 5-2(a)), while the phosphorous layer by PH3 is too thin to be clearly observed (Fig.

5-2(b)). The much lower sticking coefficient of phosphorous [8] and limitation of surface open sites for phosphorous adsorption [6] are responsible for this. It was reported that phosphorous can only reach a half of monolayer on Ge surface at saturation[9].

Ge

Si

(a)

20nm

Ge

Si

(b)

20nm

Fig.5-1 The TEM cross-sectional images of epitaxial Ge (a) as grown and (b) after anneal. Note the threading dislocation density in Ge film with post annealing is ~5x10⁸ cm^-2.

The doping layer on Ge surface is further characterized by x-ray photoelectron spectroscopy (XPS) (Fig. 5-3). Based on material database and references [10-12], the embedded peaks can be deconvoluted. For the boron layer on Ge, there are boron-interstitial-clusters (B-B) at 186.5eV and the complex of boron coordinated with Ge (B-Ge) at 187.6eV (Fig. 5-3(a)). Due to the instability of the boron-vacancy pair, the boron above the solid solubility could be precipitated through the boron-interstitial-cluster [13]. The chemical structures of such clusters in Ge are not known, but BI^-, B2I⁰, and B2I30 clusters have been found in Si [14]. For the phosphorous layer on Ge, there are phosphorous- interstitial-clusters (P-P) at 130.3eV, the complex of phosphorous coordinated with Ge (P-Ge) at 129.2eV, and Ge 3p at 126eV (Fig. 5-3(b)). The B-Ge peak and P-Ge peak are signatures of the B and P diffusion into Ge, respectively.

Ge

(a)

Ge

2 nm

(b)

Fig.5-2 The TEM cross-sectional images of (a) a ~5nm boron layer and (b) less than monolayer phosphorous layer on epitaxial Ge growth at 450^oC for 2 hrs.

Notably, the high surface doping density with sharp spatial decay, obtained by using the solid phase doping technique, is of particular advantage for the formation of abrupt dopant distribution at source/drain of scaled devices [15,16]. The boron and phosphorrous doping profiles by the secondary ion mass spectrometry (SIMS) measurement using O2+ ion and Cs⁺ ion, respectively, are shown in Fig. 5-4.The abruptness of boron profile is ~3nm/dec with the peak surface concentration of ~7x10²¹ cm^-3 (Fig. 5-4(a)) while the abruptness of phosphorous profile is ~5nm/dec , with the peak surface concentration of ~6x10¹⁹ cm^-3 ( Fig. 5-4(b) ). The target atoms are collided by the sputtering ions (O2+ or Cs⁺) can be driven into deeper depth, so called “knock-in effect”. Since the phosphorous is heavier than boron, the knock-in effect of phosphorous

189 188 187 186 185 P-P peak (130.3eV), P-Ge peak (129.2eV), and Ge 3p (126eV).

is not as significant as boron. The high boron concentration on the surface reflects the residuals of boron layers even after the top boron layers were etched. The double negatively charged vacancy (VGe-2) [17] in n-type Ge mediates the phosphorous diffusion [18] while the Ge interstitial in p-type Ge mediates the boron diffusion[19].

Because the formation energy of Ge vacancy is 1 eV lower than that of interstitial [20], the VGe-2 dominates in the epitaxial Ge layers. The phosphorous diffusion, mediated by the vacancy, has a larger roll-off slope of phosphorous (5nm/dec) than boron (3nm/dec).

Due to the calibration difficulties of the spreading resistance measurement on such thin layers, the reliable carrier profile cannot be measured.

0 20 40 60 80 100

Fig.5-4 The SIMS profiles of (a) boron dopants in Ge and (b) phosphorous dopants in Ge.

5.4 P

⁺

N or N

⁺

P Ge Diodes by Solid Phase Layer Doping

Ge p⁺/n and n⁺/p diodes are fabricated by dopant layer deposition to investigate the effects of defects on diode performance. The electrical characteristics of p⁺/n diodes doped by boron layer on epitaxial Ge has ~10⁵ on/off ratio, the reverse saturation current (Ioff) of ~1x10^-4A/cm² at -2V, and the ideality factor (n) of ~1.1 (Fig. 5-5(a) ).

The activation energy (Ea) of 0.34eV is estimated by the slope of the Arrhenius plot of the temperature-dependent reverse saturation currents (Fig. 5-5(b)).

The reverse saturation current from the Ref. 21 is

where the W is the depletion width, D is the diffusivity of carrier,is the effective lifetime, and ni is the intrinsic carrier density with an activation energy of half bandgap.

There are two parts of activation energy in our p⁺n diodes. The activation energy of the generation current, which is near the half band gap of Ge, indicates that the carrier generation from conduction band to valance band is dominant instead of the defect-assisted generation current. When the temperature increases to ~100^oC, the diffusion currents start to affect the reverse currents and the activation energy increases [21].

𝐽_𝑅 = 𝐽_{𝑑𝑖𝑓𝑓𝑢𝑠𝑖𝑜𝑛} + 𝐽𝑔𝑒𝑛𝑒𝑟𝑎𝑡𝑖𝑜𝑛 =q 𝐷 τ

𝑛_𝑖²

𝑁 +𝑞𝑊𝑛_𝑖 𝜏

89

From the Fig. 5-6 (a), the electrical characteristic of n⁺/p diodes doped by phosphorous layers on epitaxial Ge has the on/off ratio of ~1.5x10⁵, the reverse saturation current of ~4x10^-5A/cm² at -2V, and the ideality factor of ~1.2. Since there is no passivation on the sidewalls of the mesa, the ideality factor of 1.2 of n⁺/p can be the defects on the sidewalls and/or the defects near the bottom Ge on Si, where the misfit dislocations exist. The activation energy of 0.36eV (Fig. 5-6 (b)) is larger than those by ion implantation (0.15~0.26eV)[27], also indicating the carrier generation from the valance band edge to the conduction band edge.

-2 -1 0 1 2 boron layer at the room temperature. The on/off ratio is about ~1x10⁵, ideality factor is about 1.1, and the reverse saturation current is ~1x10^-4A/cm². (b) The Arrhenius plot of the reverse saturation currents with the J-V curves at different temperatures in the inset.

-2 -1 0 1 2

Fig.5-6 (a) J–V characteristics of n⁺/p Ge diodes by in-situ phosphorous layer at the room temperature. The on/off ratio is ~1.5x10⁵, ideality factor is about 1.2, and the reverse saturation current is 4x10^-5A/cm². (b) The Arrhenius plot of the reverse saturation currents with the J-V curves at different temperatures in the inset.

-2 -1 0 1 2

5.5 Summary

As compared to the reported Ge p⁺/n diodes in Table 5-I, the reverse saturation currents of our p⁺/n diodes by in situ boron layer doping is 10 to 100 times lower than that by ion implanted diodes on bulk Ge and epitaxial Ge [22-24] Free from implantation damage by in-situ boron layer doping is the root cause.

The on/off ratio also reaches the best among all the diodes. Besides, the reported n⁺/p diodes listed in the Table 5-II. Since the threading dislocations connecting from the surface to the bottom of the film are also high conductance paths which is the major cause for the leakage current in diodes, our epitaxial Ge on Si by in situ annealing after

on/off ratio I

_off

(A/cm

²

)

BF₂ion implantation

/ RTA of Ge (Ref.22)

~1x10

³

3x10

^-2

Metal induce defect

annealing of Ge (Ref.23)

~4x10

⁴

1x10

^-2

BF₂ion implantation

/RTA of epi-Ge (Ref.24)

~5x10

³

7x10

^-4

B ion implantation/ RTA on Ge

growth is also responsible for better performance than that in Ref.26. Doping Ge by the diffusion of dopant layers on Ge has much lower defect density as compared to ion implantation. The low defect density leads to abrupt dopant profiles, and good diode characteristics. This work provides an alternative path way for doping using chemical vapor deposition.

on/off ratio I

_off

(A/cm

²

)

Sb ion implantation

/ Laser on Ge (Ref.25)

~1x10

³

1x10

^-2

In situ Doping of

epi-Ge (Ref.26)

~1.1x10

⁴

1.2x10

^-2

Spin on doping on Ge

(Ref.4)

~1x10

⁶

1x10

^-4

Gas phase doping on Ge

(Ref.27)

~5x10

⁴

3x10

^-3

P ion implantation/ twice

RTA on Ge (Ref.28)

~2x10

⁵

5x10

^-5

Phosphorous Layer Doping of epi-Ge

(This work)

~1.5 x10

⁵

4x10

^-5

Table 5-2. Comparison of the n⁺/p Ge diodes by in-situ phosphorous layer doping and other methods from precious works. Our diodes shows superior Ioff [4, 25-28].

Reference

[1] Chi On Chui, L. Kulig, J. Moran, W. Tsai, and K. C. Saraswat, “ Germanium n-type shallow junction activation dependences,” Appl. Phys. Lett., Vol. 87, 091909, Aug. 2005.

[2] J. Kim, S. W. Bedell, and D. K. Sadana, “Improved germanium n⁺/p junction diodes formed by coimplantation of antimony and phosphorus,” Appl. Phys. Lett., Vol. 98,

[2] J. Kim, S. W. Bedell, and D. K. Sadana, “Improved germanium n⁺/p junction diodes formed by coimplantation of antimony and phosphorus,” Appl. Phys. Lett., Vol. 98,

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