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Characterization and strain gradient optimization of PECVD poly-SiGe layers for MEMS applications
Sensors and Actuators A 130–131 (2006) 403–410
Characterization and strain gradient optimization of PECVD poly-SiGe layers for MEMS applications M. Gromova a,b,∗ , A. Mehta a , K. Baert a , A. Witvrouw a a
b
IMEC, Leuven, Belgium Katholieke Universiteit Leuven, ESAT-INSYS, Leuven, Belgium
Received 9 June 2005; received in revised form 15 November 2005; accepted 28 December 2005 Available online 10 February 2006
Abstract Poly-SiGe offers an attractive alternative for low temperature MEMS post-processing above CMOS. This paper presents several investigations made to obtain a crystalline material with excellent mechanical (low stress, low stress gradient) and electrical (low resistivity) properties. Two different techniques were used to enhance the crystallization of the plasma enhanced chemical vapor deposition (PECVD) layers at low temperatures. The use of a high hydrogen dilution (H2 /(SiH4 + GeH4 ) ≈ 90) leads to microcrystalline SiGe (cSiGe:H) at temperatures as low as 350 ◦ C. In this work ∼0.6 and 2 m thick cSiGe:H layers were characterized and optimized. Alternatively, a chemical vapor deposition (CVD) crystallization layer can be used below the PECVD layer in order to deposit thick (4–10 m) polycrystalline films of high quality at temperatures around 450 ◦ C. For both layers, the strain gradient can be optimized by the use of a compressive top layer. In the case of the cSiGe:H layers an alternative new strain gradient optimization method, which uses a variable hydrogen dilution to effectively fine-tune the mechanical properties of the cSiGe films, is presented. Also new results on the surface roughness of the layers are presented. © 2006 Elsevier B.V. All rights reserved. Keywords: Microcrystalline; Poly-SiGe; Hydrogen dilution; Multi-layer; PECVD; Strain gradient
1. Introduction Polycrystalline-SiGe (poly-SiGe) has been demonstrated to be an ideal material for post-processing MEMS above CMOS, since films with very good electrical and mechanical properties can be obtained at CMOS-compatible temperatures [1,2]. Alloying Si with Ge lowers the melting point and the amorphous to crystalline transition temperature, and therefore also the thermal budget needed to obtain good MEMS structural layers is lowered. Using methods such as multi-layer deposition [3], a high hydrogen dilution [3,4], laser annealing and metal-induced crystallization [5] can lower the deposition temperature further (350–450 ◦ C). The first part of this work deals with the development of microcrystalline SiGe as a high quality material for MEMS applications. The theoretical background of the hydrogenated microcrystalline SiGe (cSiGe:H) deposition process was discussed in detail in [4]. A study of films of ∼0.6 and 2 m thickness is presented. To optimize the strain gradient
a compressive top layer can be used. Also a new method for strain gradient fine-tuning is introduced: by varying the hydrogen dilution ratio (H2 /(SiH4 + GeH4 )) during layer deposition it is possible to influence the structure of the SiGe crystals, resulting in a different stress profile in the film. Also with this method the strain gradient can be minimized. Furthermore, a multi-layer stack combining chemical vapor deposition (CVD) and plasma enhanced chemical vapor deposition (PECVD) depositions at 450 ◦ C is presented, resulting in films with excellent electrical and mechanical properties grown at a high deposition rate. This process is very useful for applications where thick (4–10 m) MEMS structural layers with very low strain gradients need to be deposited above standard CMOS [6]. The strain gradient of these multi-layer films was also optimized by the use of a compressive top layer. 2. Experimental 2.1. Deposition
∗
Corresponding author at: IMEC, Kapeldreef 75, B-3001 Leuven, Belgium. Tel.: +32 16 28 10 44; fax: +32 16 28 15 01. E-mail address: [email protected] (M. Gromova). 0924-4247/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.sna.2005.12.048
Crystalline SiGe layers were deposited in an Oxford plasma technology (OPT) Plasma lab 100 cold-wall PECVD system,
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which uses a standard parallel plate reactor. A gas flow of 10% germane (GeH4 ) in hydrogen (H2 ) was used as germanium (Ge) source. The silicon source was a flow of pure silane (SiH4 ). The ratio of the silane to germane (SiH4 /GeH4 ) flows determines the Ge concentration in the film. Doping of the layers is done using a 1% diborane (B2 H6 ) in H2 as boron (B) gas source (p-type SiGe). The exact doping level will influence the stress in the layers [7], but this was not investigated within this work. The layers are deposited on 6- or 8-in. silicon wafers (1 0 0) having a SiO2 layer on top. Rutherford back scattering (RBS) was used to determine the Si and Ge concentration in the layers. A four-point probe was used for the sheet resistance measurements. Morphology and grain microstructure were investigated with transmission electron microscopy (TEM). The average stress was determined by measuring the curvature of the wafer, before and after deposition of the layers, using an Eichhorn & Hausmann MX 203 stress-meter. The poly-SiGe layers were patterned and plasma etched in a deep dry etching system from surface technology systems (STS) using an SF6 + O2 /C4 F8 alternating plasma. The thickness of the layers was measured with a stylus profilometer (DekTak). Cantilevers were released after that by removing the underlying sacrificial thermal SiO2 using HF (49%) vapor at 35 ◦ C. The cantilever profile and tip deflection were measured using a Wyko profilometer from VEECO. The deflection data is used to calculate the strain gradient in the film. 2.1.1. Microcrystalline deposition process cSiGe layers are deposited at 1 Torr pressure. The plasma power used is 370 mW/cm2 . Experiments were performed at 350 and 400 ◦ C deposition temperatures. For the single films and the films with compressive top layer, a hydrogen dilution (H2 /(SiH4 + GeH4 )) of approximately 90 was used. The H2 and B2 H6 flows were kept constant throughout the deposition at 2 slm and 10 sccm, respectively, while the SiH4 /GeH4 was variable, determining the Ge content in the grown film. For the variable hydrogen dilution depositions (see Section 3.1.2.2) H2 /(SiH4 + GeH4 ) ratios as high as 90 (2 slm H2 flow)
were used at the start of the deposition, followed by steps with a gradually decreasing hydrogen flow. 2.1.2. Multi-layer combining CVD and PECVD The multi-layers are deposited at 2 Torr pressure, 450 ◦ C and 0 W power (no plasma) for the (LP)CVD part of the layer. The PECVD parts of the layer are normally deposited at 30 W (power density 61 mW/cm2 ) except for the compressive top layer deposited at high plasma power (40 W). No extra hydrogen dilution was used for these films. A constant B2 H6 flow equal to 40 sccm was used during the deposition of the entire layer. A SiH4 /GeH4 ratio of 2.4 was used for the whole layer stack except for the Si-rich top compensation layer and the a-Si seed layer. 3. Results and discussion 3.1. Microcrystalline deposition process 3.1.1. Single film characterization Fig. 1 shows the TEM cross-sections of as-grown thin (∼0.6 m thick) single microcrystalline SiGe layers deposited at 350 ◦ C with different Ge content. The pure Ge layer is completely amorphous (Fig. 1a), while the pure Si layer is fully crystalline (Fig. 1b) and consists of fine grains. SiGe layers have well defined almost columnar grains and a small amorphous fraction at the SiO2 /SiGe interface (Fig. 1c). The surface roughness of these layers was measured using an atomic force microscope (AFM), and the root-mean-square (RMS) values are presented in Fig. 2. The surface roughness gradually increases with increasing Ge content in the layer. For example, the RMS value is 4.6 nm for a film with 44% Ge, while the RMS value is 7.3 nm in the case of a film with 69% Ge. It is worth noticing the two extreme cases: a pure Si layer has a roughness of 5.7 nm, while the pure Ge layer has a roughness of only 4 nm. This is to be expected since the pure Ge layer is completely amorphous, as proven by the TEM image (Fig. 1a). This means that the use of the high hydrogen dilution is much more efficient in obtaining crystalline layers at low deposition temperatures in the case of SiGe and Si than in the case of pure Ge, as also
Fig. 1. Cross-section TEM of as-deposited single layers deposited at 350 ◦ C—(a) pure Ge: completely amorphous; (b) pure Si: fully crystalline, consisting of fine grains; (c) cSi34 Ge66 : well defined columnar grains, small amorphous fraction at the SiO2 /SiGe interface.
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Table 1 Properties of ∼0.6 m thick Si31 Ge69 layer boron doped cSiGe:H deposited at 350 ◦ C, 1 Torr, 370 mW/cm2 power density with different top layers
Fig. 2. Surface roughness of ∼0.6 m thick cSiGe films deposited at 350 ◦ C as a function of Ge content. The pure Ge film has a very smooth surface since it is amorphous.
Fig. 3. Stress as a function of Ge content in ∼0.6 m thick cSiGe films deposited at 350 ◦ C.
shown in literatures [8,9]. Also the stress as a function of the Ge content was determined for layers deposited at 350 ◦ C (Fig. 3). It is obvious that the stress changes from being highly compressive in layers with high Si content to low tensile in layers with a high Ge content. As the stress gradient in single layers normally causes released cantilevers to bend upwards (see Section 3.1.2.1), a beneficial effect on the stress and stress gradient is expected from combining SiGe layers with high Ge content with a Si-rich layer (Fig. 4) [10]. Thus, a 0.5 m thick Si31 Ge69
Layer
Thickness (m)
Resistivity (m cm)
Stress (MPa)
Si31 Ge69 + 0.1 m Si34 Ge66 Si31 Ge69 + 0.1 m Si40 Ge60 Si31 Ge69 + 0.1 m Si47 Ge53
0.64 0.6 0.6
23 22 23
+29 +17 +8
layer (+25 MPa) was combined with an 0.1 m thick ‘Si-rich’ layer (compressive stress). The results, presented in Table 1, prove that the bi-layers have a low resistivity and a low tensile stress (between +29 and +8 MPa, depending on the top layer used). The strain gradient behavior of these layers still has to be studied, but we are confident through our experiments shown in Section 3.1.2.1, that by combining layers of the right thickness and appropriate Ge concentrations, thin layers with near-zero strain gradients can be grown in situ. 3.1.2. Strain gradient optimization 3.1.2.1. Use of compressive top layer. Similar experiments as described in Section 3.1.1 were performed for layers as thick as 2 m. First, a number of experiments were done to determine the stress and strain gradient for films with different Ge content at two deposition temperatures (Table 2). Similar as for the thin layers, the stress of the cSiGe:H layers deposited both at 350 and 400 ◦ C varies from compressive to tensile with increasing Ge content. Therefore, also for these thicker layers it is expected that the strain gradient can be fine-tuned by combining layers with different Ge content. For example, combining a 2 m thick Si33 Ge67 layer, with a 0.3 m thick Si52 Ge48 top layer resulted in a stress of +32 MPa, a resistivity of 12 m cm and a strain gradient of 1.5 × 10−4 m−1 for deposition at 400 ◦ C (Table 3), while a single layer deposited using similar conditions has a strain gradient of 3 × 10−4 m−1 . The deposited bi-layer is crystalline as shown in Fig. 5. The same stack deposited at 350 ◦ C leads to a
Fig. 4. Stress gradient optimization by stacking layers of opposing stress behavior.
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Table 2 Properties of boron doped single cSiGe:H layers deposited at 1 Torr, 370 mW/cm2 power density and H2 /(SiH4 + GeH4 ) ≈ 90, layer thickness ∼2 m Temperature (◦ C)
Ge (%)
Deposition rate (nm/min)
Resistivity (m cm)
Stress (MPa)
Strain gradient (×10−4 m−1 )
350 350 350 400 400
56 45 42 67 48
23 21 17 23 20
15 14 18 9 11
+48 −135 −238 +25 −32
13.9 5 14 3 6.3
Table 3 Properties of 2 m thick Si33 Ge67 layer boron doped cSiGe:H deposited at 1 Torr, 370 mW/cm2 power density and a top layer (Si52 Ge48 ) of different thickness Temperature (◦ C)
Si-rich layer thickness (m)
Resistivity (m cm)
Stress (MPa)
Strain gradient (×10−4 m−1 )
400 400 400 400 350
3 1 0.6 0.3 0.3
8 11 12 12 13
+46 +28 +34 +32 −75
Bend down Bend down − 1.5 4
compressive stress of 75 MPa, a resistivity of 13 m cm and a strain gradient of 4 × 10−4 m−1 (Table 3). A single layer grown at 350 ◦ C has a strain gradient as high as 14 × 10−4 m−1 . 3.1.2.2. Use of variable H2 dilution. Hydrogen plays a key role in obtaining crystalline material at low temperatures and in reducing defects during growth. The results presented above prove that the high hydrogen dilution used during deposition is an effective method for obtaining high quality crystalline
Fig. 5. Cross-section TEM of as-deposited bi-layer cSiGe (bottom layer: Si33 Ge67 ; top layer: Si52 Ge48 ) deposited at 400 ◦ C.
SiGe layers at temperatures as low as 350–400 ◦ C. Here, a novel method for improving the crystalline structure and optimization of the mechanical properties of the film is proposed, consisting in a gradual reduction of the hydrogen dilution during the deposition of the film. As the H2 dilution helps the crystallization process, decreasing the H2 dilution is expected to lead to smaller crystals and/or an increased amorphous fraction in the layers grown with the reduced dilution. A gradual decrease in H2 dilution is thus expected to lead to a gradual decrease in tensile stress with film thickness and will thus give a means to compensate the positive strain gradient of the single films grown with constant H2 dilution. Actually, a reduction in H2 dilution might even lead to more columnar grain growth than in the case of single films as the positive strain gradient in single films is probably caused not only by the amorphous fraction often seen at the interface of the film, but also by some increase in grain size over the film thickness. An additional benefit of this method of strain gradient optimization is that no variation in Ge content over the layer is expected (in contrast to the method where a top Si-rich layer is used), which can be important for the thermal stability of the SiGe films. Also, the overall growth rate of the deposition process increases with reduced H2 dilution (see below). An overview of the performed experiments and the characteristics of the films are reported in Table 4. Two different conditions were used as a basis: (A) SiH4 /GeH4 = 1.8 and (B) SiH4 /GeH4 = 1.3. Within these two base conditions the H2 dilution was reduced in different steps. The first step is performed using a H2 flow of 2 slm, leading to a dilution as high as 90 (condition A) or 76 (condition B). This high dilution is necessary in order to assure the initial nucleation of the layer. In the same time, it is known [11,12] that H2 etches away the less energetically favorable bonds from the film surface, promoting so the growth of a higher quality crystalline material, reducing in the same time the deposition rate. Therefore, reducing the H2 dilution for each step compared to the previous one might have a two-fold beneficial influence on the growth of the film: an increase of the deposition rate and the growth of more columnar grains instead of V-shaped ones. TEM analysis was done to acquire the necessary information on the microstructure. The grain microstructure through part of these layers is shown for condition A1 in Fig. 6 and for condition A3 in Fig. 7. A small amorphous fraction at the SiO2 /SiGe is observed in both cases. V-shaped grains are found at the beginning of the layer, which transform into elongated columnar grains within the thickness of the film. In [13] it was observed that during CVD growth of SiGe V-shaped grains are formed, as many grains
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Table 4 Properties of ‘thin’ (<1 m) boron doped cSiGe:H deposited at 350 ◦ C, 1 Torr, 370 mW/cm2 power density and hydrogen dilution (H2 /(SiH4 + GeH4 )) decreasing during deposition SiH4 /GeH4
Condition
H2 dilution (H2 /(SiH4 + GeH4 ))
Time in step (min)
Thickness (m)
Deposition rate (nm/min)
1.8
A1
90 68 46.5 25
10 5 5 5
0.7
28
32
−35
A2
90 46.5
10 10
0.53
26.5
27
−45
A3
90 79 68 46.5
10 5 5 5
0.64
25.6
21.8
−32
B1
76 61 42 23
5 5 5 5
0.7
34
59
+29
B2
76 42 23 14
5 5 5 5
0.8
39.5
274
−19
B3
76 70 61 42
5 5 5 5
0.6
29.5
48
+53
1.3
Resistivity (m cm)
Stress (MPa)
Condition B1: strain gradient = 3.6 × 10−5 m−1 .
start to grow, but only a few reach the top. This results in a wide variation in grain size over the film thickness leading to a stress gradient through the layer. For the faster growing PECVD films [13] more columnar-like grains and thus a reduced stress gradi-
ent was observed compared to CVD SiGe for films with equal thickness. Based on this, we believe that the very high hydrogen dilution should be used in the beginning of the growth process, where nucleation has to be achieved and V-shaped grains start to grow. Afterwards, a reduced H2 dilution can be used for ‘tuning’ the shape of the grains, restricting them to grow more columnar-
Fig. 6. Cross-section TEM of 0.84 m thick cSiGe deposited using variable H2 dilution, Table 4, condition A1. Elongated columnar grains are observed, smaller grains at the SiGe-oxide interface, partially amorphous fraction.
Fig. 7. Cross-section TEM of 0.74 m thick cSiGe deposited using variable H2 dilution, Table 4, condition A3 (top of the picture: substrate side; bottom: film surface). Elongated columnar grains are observed.
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like and therefore giving us a possibility to tune the strain gradient. For condition B1 (Table 4) a very low strain gradient of 3.6 × 10−5 m−1 was measured, which is a factor of 10 better than the best value obtained using a compressive top layer at 350 ◦ C. 3.2. Multi-layer combining CVD and PECVD 3.2.1. Film characterization The microcrystalline deposition method results in excellent films at very low temperatures. Due to the hydrogen dilution however, a rather limited growth rate is obtained. Even in the case of the variable hydrogen dilution, the growth rate stays below 40 nm/min. Sometimes, e.g. for inertial sensors, the deposition of very thick layers (up to 10 m) is needed. The use of PECVD without hydrogen dilution would in that case be more economical. However, PECVD depositions alone at CMOS-compatible temperatures (∼450 ◦ C or lower) give amorphous films characterized by a high compressive stress and high resistivity values [3]. An LPCVD process gives crystalline films at these low temperatures, but has a very low deposition rate. Hence, it was shown in [3] that combining the two processes gives the benefits of both PECVD and CVD (without plasma), namely a high deposition rate and a low crystallization temperature, respectively. The polycrystalline SiGe multi-layer stack used in this work is composed of a thin PECVD seed layer (a-Si or a-SiGe), a LPCVD SiGe crystallization layer and a bulk PECVD SiGe layer (Fig. 8). The process time for the last PECVD step determines the
Fig. 8. Cross-section TEM for 1 m multi-layer film showing columnar grains (Si35 Ge65 ).
Fig. 9. Deflection of a 1 mm long released cantilever for 10 m thick SiGe layer with 800 nm Si-rich top layer.
thickness of the layer and the average deposition rate. While this process was introduced in a previous paper [3], in this work the strain gradient is further fine-tuned (down to 3.5 × 10−6 m−1 ) using a top compressive layer. A 1 m film exhibits a low average stress of −5 MPa, and an average resistivity of 1 m cm. It has a polycrystalline structure with columnar grains. Ten micrometers thick layers exhibit high deposition rates (∼100 nm/min), a low resistivity of 0.9 m cm, a tensile stress of 71 MPa and a positive strain gradient of 3.6 × 10−5 m−1 . 3.2.2. Strain gradient optimization Similar as shown for the microcrystalline films, it is possible to modify the stress and strain gradient in these multi-layers by stacking layers of opposing stress behavior. Thus, a top compensation layer on the multi-layer stack could further reduce the strain gradient. A few options that were explored for a 10 m thick film with a-SiGe seed layer are as follows: (a) a CVD top layer, (b) a top layer deposited at high plasma power and (c) a silicon rich top layer. Of these, a 10 m film with 1 m silicon-rich top layer gives the best results with a tensile stress of 57 MPa and a strain gradient of 1.2 × 10−5 m−1 [3]. The thickness and composition of the top compensation layer depends on the total
Fig. 10. Stress profile of a 11.6 m thick layer rebuilt from stress gradient measurements done after successive thinning down the layers using low energy plasma.
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Fig. 11. Thick poly-SiGe gyroscope on top of standard 0.35 m Al-CMOS process.
film stack built up (including the seed layer and the underlying sacrificial layer) and the stress gradient aimed for, as shown below. For example, a 10 m thick film with an a-Si seed layer and a 800 nm thick Si-rich compensation layer exhibited an average stress of 35 MPa, a resistivity of 1.45 m cm, a deposition rate of 90 nm/min and an extremely low strain gradient value (3.6 × 10−6 m−1 ) (Fig. 9). The amorphous seed layer is highly compressive and is expected to have a big influence on the overall strain gradient. This is apparent from the fact that a much thicker (1 m) top compensation layer is required to achieve a strain gradient of 1.2 × 10−5 m−1 for an a-SiGe seed layer (see above) as compared to an a-Si seed layer, where a 800 nm thick compensation layer gives a strain gradient of 3.6 × 10−6 m−1 . For a similar 10 m thick poly-SiGe layer, the stress profile across its thickness was derived. This was achieved by successively thinning down the layer, using very low energy plasma to avoid damaging the film, and measuring the strain gradient at each thickness. The strain gradient data was then used to construct the complete stress profile [14]. Fig. 10 shows that the maximum stress variation is in the initial layers after which the stress becomes more or less constant. The top compensation layer reverses the stress profile. Thus, at first it was expected that, while keeping the seed layer and top compensation layer constant, and just by varying the thickness of the bulk PECVD layer, it should be possible to obtain similar strain gradient values for different film thickness, with very little optimization required. Starting with a similar stack as for the 10 m layer, 4 m thick poly-SiGe films were developed. In this case however, the compensation layer needed to be thicker than expected to achieve
a low strain gradient. For a 1.35 m thick top layer, a stress value of 20 MPa, a resistivity of 1 m cm and a strain gradient of 3.5 × 10−6 m−1 is obtained. This result shows that, although the top compensation layer needed fine-tuning, it is possible to realize low strain gradient layers at different thickness. The optimized films described here can be used as structural layers for CMOS-integrated gyroscopes (Fig. 11) [15], accelerometers, etc. As these multi-layer films with top compressive layer have a variable Ge content over their layer thickness, the stability of these layers with time and temperature initially was a concern. In [6] it is however shown that both 4 and 10 m thick films are very stable over time and during temperature variations between room temperature and 100 ◦ C. 4. Conclusions The work presented above gives ample evidence that both microcrystalline SiGe and CVD/PECVD multi-layer poly-SiGe films form excellent MEMS structural layers, especially of interest for low thermal budget applications. Excellent electrical and mechanical properties were obtained with extremely low strain gradients (<2 m deflection for 1 mm long cantilever). These low strain gradients can be either reached through the use of top compressive layers for both types of films or through the use of a variable hydrogen dilution in the case of microcrystalline SiGe. These films are also stable with respect to time and temperature. First CMOS-integrated moving gyroscopes have been realized and other applications are currently being evaluated.
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Acknowledgements The work on 10 m thick films presented here is partially supported by the European project “SiGeM” (IST-2001-37681). The authors would also like to acknowledge the SiGeM project partners and Tim Brosnihan for the valuable discussions and suggestions. References [1] A. Witvrouw, M. Gromova, A. Mehta, S. Sedky, P. De Moor, K. Baert, C. Van Hoof, Poly-SiGe, a superb material for MEMS, Proc. MRS 782 (2004) 25–36. [2] R.T. Howe, T.J. King, Low-temperature LPCVD MEMS technologies, MRS Spring Symp. U 729 (2002) 205–213. [3] A. Mehta, M. Gromova, C. Rusu, O. Richard, K. Baert, C. Van Hoof, A. Witvrouw, Novel high growth rate processes for depositing poly-SiGe structural layers at CMOS compatible temperatures, in: Proceedings of the MEMS’04, 2004, pp. 721–724. [4] M. Gromova, K. Baert, C. Van Hoof, A. Mehta, A. Witvrouw, The novel use of low temperature hydrogenated microcrystalline silicon germanium for MEMS applications, Microelectron. Eng. 76 (2004) 266– 271. [5] T.J. King, R.T. Howe, S. Sedky, G. Liu, B.C. Lin, M. Wasilik, C. Duenn, Recent progress in modularly integrated MEMS technologies, Digest IEEE IEDM (2002) 199–202. [6] A. Mehta, M. Gromova, P. Czarnecki, K. Baert, A. Witvrouw, Optimization of PECVD poly-SiGe layers for MEMS post-processing on top of CMOS, in: Proceedings of the Transducers’05, 2005, pp. 1326– 1329. [7] S. Sedky, A. Witvrouw, A. Saerens, P. Van Houtte, J. Portmans, K. Baert, Effect of in situ boron doping on properties of SiGe films deposited by chemical vapor deposition at 400 ◦ C, J. Mater. Res. 16 (9) (2001) 2607–2612. [8] H.C. Lin, C.Y. Chang, W.H. Chen, W.C. Tsai, T.C. Chang, T.G. Jung, H.Y. Lin, Effects of SiH4 , GeH4 and B2 H6 on the nucleation and deposition of polycrystalline Si1−x Gex films, J. Electrochem. Soc. 141 (9) (1994) 2559–2563. [9] G. Ganguly, T. Ikeda, T. Nishimiya, K. Saitoh, M. Kondo, A. Matsuda, Hydrogenated microcrystalline silicon germanium: a bottom cell material for amourphous silicon-based tandem solar cells, Appl. Phys. Lett. 69 (27) (1996) 4224–4226. [10] J. Yang, H. Kahn, A.-Q. He, S.M. Phillips, A.H. Heuer, A new technique for producing large-area as-deposited zero-stress LPCVD polysilicon films: the multipoly process, J. Microelectromech. Syst. 9 (4) (2000) 485–494. [11] K. Tanaka, A. Matsuda, Glow Discharge Amorphous Si Growth Process and Structure, NHPC, Amsterdam, 1987, pp. 156–158. [12] K. Baert, Very Low Temperature Growth of Microcrystalline and Epitaxial Silicon by PECVD, KULeuven, Leuven, 1990, pp. 9– 10.
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Biographies Maria Gromova received an MS degree in Communication Engineering from the Technical University of Sofia, Bulgaria in 1999. She is currently pursuing a PhD degree in Electronics Engineering at IMEC, Leuven, with main research interests including post-CMOS integration of MEMS. In particular hydrogenated microcrystalline SiGe development for MEMS applications. Anshu Mehta received her Bachelor’s degree from Indian Institute of Technology, Bombay in 2000, and a MS in engineering science from Pennsylvania State University in 2002. She has been working at Interuniversity MicroElectronics Center, Leuven, as a Research Engineer, since 2002. Her research interests include poly-SiGe surface micromachining, focusing on CMOSintegrated MEMS. Kris Baert received his PhD degree in microelectronics and materials science at Leuven University, Belgium in 1990. His dissertation was on lowtemperature plasma-enhanced CVD of amorphous and microcrystalline Si for solar cells. In 1988, he was a Visiting Scientist at the Konagai-Takahashi Laboratory, Tokio Institute of Technology, Japan. From 1990 to 1992, he was a Research Scientist at the Materials and Electronics Devices Laboratory, Mitsubishi Electric, Amagasaki, Japan, developing poly-Si TFTs. In 1992, he joined the Interuniversity Micro-electronics Center, Leuven, as a Senior Scientist working on micromachined IR sensors. Since 1995, he has been responsible for development of microsystem technologies. He is currently head of the MEMS group working on RF-MEMS, poly-SiGe surface micromachining and CMOS integration. Ann Witvrouw received the MS degree in metallurgical engineering in 1986 from the Katholieke Universiteit, Leuven, Belgium, and both the MS degree in applied physics in 1987 and PhD degree in applied physics in 1992 from Harvard University, Cambridge, MA. In 1992, she joined the Interuniversity Microelectronics Center (IMEC), Leuven, where she worked on the reliability of metal interconnects until the end of 1998. During this time, her research was focused on the mechanical stress in films and lines, electromigration and stress induced voiding. In 1998, she switched to research in Micro-electromechanical Systems at IMEC, where she is now responsible for advanced MEMS process technologies including the integration of MEMS and CMOS. She has been the coordinator of the IST project SUMICAP and currently she is the coordinator of the IST project SiGeM.
Characterization and strain gradient optimization of PECVD poly-SiGe layers for MEMS applications M. Gromova a,b,∗ , A. Mehta a , K. Baert a , A. Witvrouw a a
b
IMEC, Leuven, Belgium Katholieke Universiteit Leuven, ESAT-INSYS, Leuven, Belgium
Received 9 June 2005; received in revised form 15 November 2005; accepted 28 December 2005 Available online 10 February 2006
Abstract Poly-SiGe offers an attractive alternative for low temperature MEMS post-processing above CMOS. This paper presents several investigations made to obtain a crystalline material with excellent mechanical (low stress, low stress gradient) and electrical (low resistivity) properties. Two different techniques were used to enhance the crystallization of the plasma enhanced chemical vapor deposition (PECVD) layers at low temperatures. The use of a high hydrogen dilution (H2 /(SiH4 + GeH4 ) ≈ 90) leads to microcrystalline SiGe (cSiGe:H) at temperatures as low as 350 ◦ C. In this work ∼0.6 and 2 m thick cSiGe:H layers were characterized and optimized. Alternatively, a chemical vapor deposition (CVD) crystallization layer can be used below the PECVD layer in order to deposit thick (4–10 m) polycrystalline films of high quality at temperatures around 450 ◦ C. For both layers, the strain gradient can be optimized by the use of a compressive top layer. In the case of the cSiGe:H layers an alternative new strain gradient optimization method, which uses a variable hydrogen dilution to effectively fine-tune the mechanical properties of the cSiGe films, is presented. Also new results on the surface roughness of the layers are presented. © 2006 Elsevier B.V. All rights reserved. Keywords: Microcrystalline; Poly-SiGe; Hydrogen dilution; Multi-layer; PECVD; Strain gradient
1. Introduction Polycrystalline-SiGe (poly-SiGe) has been demonstrated to be an ideal material for post-processing MEMS above CMOS, since films with very good electrical and mechanical properties can be obtained at CMOS-compatible temperatures [1,2]. Alloying Si with Ge lowers the melting point and the amorphous to crystalline transition temperature, and therefore also the thermal budget needed to obtain good MEMS structural layers is lowered. Using methods such as multi-layer deposition [3], a high hydrogen dilution [3,4], laser annealing and metal-induced crystallization [5] can lower the deposition temperature further (350–450 ◦ C). The first part of this work deals with the development of microcrystalline SiGe as a high quality material for MEMS applications. The theoretical background of the hydrogenated microcrystalline SiGe (cSiGe:H) deposition process was discussed in detail in [4]. A study of films of ∼0.6 and 2 m thickness is presented. To optimize the strain gradient
a compressive top layer can be used. Also a new method for strain gradient fine-tuning is introduced: by varying the hydrogen dilution ratio (H2 /(SiH4 + GeH4 )) during layer deposition it is possible to influence the structure of the SiGe crystals, resulting in a different stress profile in the film. Also with this method the strain gradient can be minimized. Furthermore, a multi-layer stack combining chemical vapor deposition (CVD) and plasma enhanced chemical vapor deposition (PECVD) depositions at 450 ◦ C is presented, resulting in films with excellent electrical and mechanical properties grown at a high deposition rate. This process is very useful for applications where thick (4–10 m) MEMS structural layers with very low strain gradients need to be deposited above standard CMOS [6]. The strain gradient of these multi-layer films was also optimized by the use of a compressive top layer. 2. Experimental 2.1. Deposition
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Corresponding author at: IMEC, Kapeldreef 75, B-3001 Leuven, Belgium. Tel.: +32 16 28 10 44; fax: +32 16 28 15 01. E-mail address: [email protected] (M. Gromova). 0924-4247/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.sna.2005.12.048
Crystalline SiGe layers were deposited in an Oxford plasma technology (OPT) Plasma lab 100 cold-wall PECVD system,
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which uses a standard parallel plate reactor. A gas flow of 10% germane (GeH4 ) in hydrogen (H2 ) was used as germanium (Ge) source. The silicon source was a flow of pure silane (SiH4 ). The ratio of the silane to germane (SiH4 /GeH4 ) flows determines the Ge concentration in the film. Doping of the layers is done using a 1% diborane (B2 H6 ) in H2 as boron (B) gas source (p-type SiGe). The exact doping level will influence the stress in the layers [7], but this was not investigated within this work. The layers are deposited on 6- or 8-in. silicon wafers (1 0 0) having a SiO2 layer on top. Rutherford back scattering (RBS) was used to determine the Si and Ge concentration in the layers. A four-point probe was used for the sheet resistance measurements. Morphology and grain microstructure were investigated with transmission electron microscopy (TEM). The average stress was determined by measuring the curvature of the wafer, before and after deposition of the layers, using an Eichhorn & Hausmann MX 203 stress-meter. The poly-SiGe layers were patterned and plasma etched in a deep dry etching system from surface technology systems (STS) using an SF6 + O2 /C4 F8 alternating plasma. The thickness of the layers was measured with a stylus profilometer (DekTak). Cantilevers were released after that by removing the underlying sacrificial thermal SiO2 using HF (49%) vapor at 35 ◦ C. The cantilever profile and tip deflection were measured using a Wyko profilometer from VEECO. The deflection data is used to calculate the strain gradient in the film. 2.1.1. Microcrystalline deposition process cSiGe layers are deposited at 1 Torr pressure. The plasma power used is 370 mW/cm2 . Experiments were performed at 350 and 400 ◦ C deposition temperatures. For the single films and the films with compressive top layer, a hydrogen dilution (H2 /(SiH4 + GeH4 )) of approximately 90 was used. The H2 and B2 H6 flows were kept constant throughout the deposition at 2 slm and 10 sccm, respectively, while the SiH4 /GeH4 was variable, determining the Ge content in the grown film. For the variable hydrogen dilution depositions (see Section 3.1.2.2) H2 /(SiH4 + GeH4 ) ratios as high as 90 (2 slm H2 flow)
were used at the start of the deposition, followed by steps with a gradually decreasing hydrogen flow. 2.1.2. Multi-layer combining CVD and PECVD The multi-layers are deposited at 2 Torr pressure, 450 ◦ C and 0 W power (no plasma) for the (LP)CVD part of the layer. The PECVD parts of the layer are normally deposited at 30 W (power density 61 mW/cm2 ) except for the compressive top layer deposited at high plasma power (40 W). No extra hydrogen dilution was used for these films. A constant B2 H6 flow equal to 40 sccm was used during the deposition of the entire layer. A SiH4 /GeH4 ratio of 2.4 was used for the whole layer stack except for the Si-rich top compensation layer and the a-Si seed layer. 3. Results and discussion 3.1. Microcrystalline deposition process 3.1.1. Single film characterization Fig. 1 shows the TEM cross-sections of as-grown thin (∼0.6 m thick) single microcrystalline SiGe layers deposited at 350 ◦ C with different Ge content. The pure Ge layer is completely amorphous (Fig. 1a), while the pure Si layer is fully crystalline (Fig. 1b) and consists of fine grains. SiGe layers have well defined almost columnar grains and a small amorphous fraction at the SiO2 /SiGe interface (Fig. 1c). The surface roughness of these layers was measured using an atomic force microscope (AFM), and the root-mean-square (RMS) values are presented in Fig. 2. The surface roughness gradually increases with increasing Ge content in the layer. For example, the RMS value is 4.6 nm for a film with 44% Ge, while the RMS value is 7.3 nm in the case of a film with 69% Ge. It is worth noticing the two extreme cases: a pure Si layer has a roughness of 5.7 nm, while the pure Ge layer has a roughness of only 4 nm. This is to be expected since the pure Ge layer is completely amorphous, as proven by the TEM image (Fig. 1a). This means that the use of the high hydrogen dilution is much more efficient in obtaining crystalline layers at low deposition temperatures in the case of SiGe and Si than in the case of pure Ge, as also
Fig. 1. Cross-section TEM of as-deposited single layers deposited at 350 ◦ C—(a) pure Ge: completely amorphous; (b) pure Si: fully crystalline, consisting of fine grains; (c) cSi34 Ge66 : well defined columnar grains, small amorphous fraction at the SiO2 /SiGe interface.
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Table 1 Properties of ∼0.6 m thick Si31 Ge69 layer boron doped cSiGe:H deposited at 350 ◦ C, 1 Torr, 370 mW/cm2 power density with different top layers
Fig. 2. Surface roughness of ∼0.6 m thick cSiGe films deposited at 350 ◦ C as a function of Ge content. The pure Ge film has a very smooth surface since it is amorphous.
Fig. 3. Stress as a function of Ge content in ∼0.6 m thick cSiGe films deposited at 350 ◦ C.
shown in literatures [8,9]. Also the stress as a function of the Ge content was determined for layers deposited at 350 ◦ C (Fig. 3). It is obvious that the stress changes from being highly compressive in layers with high Si content to low tensile in layers with a high Ge content. As the stress gradient in single layers normally causes released cantilevers to bend upwards (see Section 3.1.2.1), a beneficial effect on the stress and stress gradient is expected from combining SiGe layers with high Ge content with a Si-rich layer (Fig. 4) [10]. Thus, a 0.5 m thick Si31 Ge69
Layer
Thickness (m)
Resistivity (m cm)
Stress (MPa)
Si31 Ge69 + 0.1 m Si34 Ge66 Si31 Ge69 + 0.1 m Si40 Ge60 Si31 Ge69 + 0.1 m Si47 Ge53
0.64 0.6 0.6
23 22 23
+29 +17 +8
layer (+25 MPa) was combined with an 0.1 m thick ‘Si-rich’ layer (compressive stress). The results, presented in Table 1, prove that the bi-layers have a low resistivity and a low tensile stress (between +29 and +8 MPa, depending on the top layer used). The strain gradient behavior of these layers still has to be studied, but we are confident through our experiments shown in Section 3.1.2.1, that by combining layers of the right thickness and appropriate Ge concentrations, thin layers with near-zero strain gradients can be grown in situ. 3.1.2. Strain gradient optimization 3.1.2.1. Use of compressive top layer. Similar experiments as described in Section 3.1.1 were performed for layers as thick as 2 m. First, a number of experiments were done to determine the stress and strain gradient for films with different Ge content at two deposition temperatures (Table 2). Similar as for the thin layers, the stress of the cSiGe:H layers deposited both at 350 and 400 ◦ C varies from compressive to tensile with increasing Ge content. Therefore, also for these thicker layers it is expected that the strain gradient can be fine-tuned by combining layers with different Ge content. For example, combining a 2 m thick Si33 Ge67 layer, with a 0.3 m thick Si52 Ge48 top layer resulted in a stress of +32 MPa, a resistivity of 12 m cm and a strain gradient of 1.5 × 10−4 m−1 for deposition at 400 ◦ C (Table 3), while a single layer deposited using similar conditions has a strain gradient of 3 × 10−4 m−1 . The deposited bi-layer is crystalline as shown in Fig. 5. The same stack deposited at 350 ◦ C leads to a
Fig. 4. Stress gradient optimization by stacking layers of opposing stress behavior.
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Table 2 Properties of boron doped single cSiGe:H layers deposited at 1 Torr, 370 mW/cm2 power density and H2 /(SiH4 + GeH4 ) ≈ 90, layer thickness ∼2 m Temperature (◦ C)
Ge (%)
Deposition rate (nm/min)
Resistivity (m cm)
Stress (MPa)
Strain gradient (×10−4 m−1 )
350 350 350 400 400
56 45 42 67 48
23 21 17 23 20
15 14 18 9 11
+48 −135 −238 +25 −32
13.9 5 14 3 6.3
Table 3 Properties of 2 m thick Si33 Ge67 layer boron doped cSiGe:H deposited at 1 Torr, 370 mW/cm2 power density and a top layer (Si52 Ge48 ) of different thickness Temperature (◦ C)
Si-rich layer thickness (m)
Resistivity (m cm)
Stress (MPa)
Strain gradient (×10−4 m−1 )
400 400 400 400 350
3 1 0.6 0.3 0.3
8 11 12 12 13
+46 +28 +34 +32 −75
Bend down Bend down − 1.5 4
compressive stress of 75 MPa, a resistivity of 13 m cm and a strain gradient of 4 × 10−4 m−1 (Table 3). A single layer grown at 350 ◦ C has a strain gradient as high as 14 × 10−4 m−1 . 3.1.2.2. Use of variable H2 dilution. Hydrogen plays a key role in obtaining crystalline material at low temperatures and in reducing defects during growth. The results presented above prove that the high hydrogen dilution used during deposition is an effective method for obtaining high quality crystalline
Fig. 5. Cross-section TEM of as-deposited bi-layer cSiGe (bottom layer: Si33 Ge67 ; top layer: Si52 Ge48 ) deposited at 400 ◦ C.
SiGe layers at temperatures as low as 350–400 ◦ C. Here, a novel method for improving the crystalline structure and optimization of the mechanical properties of the film is proposed, consisting in a gradual reduction of the hydrogen dilution during the deposition of the film. As the H2 dilution helps the crystallization process, decreasing the H2 dilution is expected to lead to smaller crystals and/or an increased amorphous fraction in the layers grown with the reduced dilution. A gradual decrease in H2 dilution is thus expected to lead to a gradual decrease in tensile stress with film thickness and will thus give a means to compensate the positive strain gradient of the single films grown with constant H2 dilution. Actually, a reduction in H2 dilution might even lead to more columnar grain growth than in the case of single films as the positive strain gradient in single films is probably caused not only by the amorphous fraction often seen at the interface of the film, but also by some increase in grain size over the film thickness. An additional benefit of this method of strain gradient optimization is that no variation in Ge content over the layer is expected (in contrast to the method where a top Si-rich layer is used), which can be important for the thermal stability of the SiGe films. Also, the overall growth rate of the deposition process increases with reduced H2 dilution (see below). An overview of the performed experiments and the characteristics of the films are reported in Table 4. Two different conditions were used as a basis: (A) SiH4 /GeH4 = 1.8 and (B) SiH4 /GeH4 = 1.3. Within these two base conditions the H2 dilution was reduced in different steps. The first step is performed using a H2 flow of 2 slm, leading to a dilution as high as 90 (condition A) or 76 (condition B). This high dilution is necessary in order to assure the initial nucleation of the layer. In the same time, it is known [11,12] that H2 etches away the less energetically favorable bonds from the film surface, promoting so the growth of a higher quality crystalline material, reducing in the same time the deposition rate. Therefore, reducing the H2 dilution for each step compared to the previous one might have a two-fold beneficial influence on the growth of the film: an increase of the deposition rate and the growth of more columnar grains instead of V-shaped ones. TEM analysis was done to acquire the necessary information on the microstructure. The grain microstructure through part of these layers is shown for condition A1 in Fig. 6 and for condition A3 in Fig. 7. A small amorphous fraction at the SiO2 /SiGe is observed in both cases. V-shaped grains are found at the beginning of the layer, which transform into elongated columnar grains within the thickness of the film. In [13] it was observed that during CVD growth of SiGe V-shaped grains are formed, as many grains
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Table 4 Properties of ‘thin’ (<1 m) boron doped cSiGe:H deposited at 350 ◦ C, 1 Torr, 370 mW/cm2 power density and hydrogen dilution (H2 /(SiH4 + GeH4 )) decreasing during deposition SiH4 /GeH4
Condition
H2 dilution (H2 /(SiH4 + GeH4 ))
Time in step (min)
Thickness (m)
Deposition rate (nm/min)
1.8
A1
90 68 46.5 25
10 5 5 5
0.7
28
32
−35
A2
90 46.5
10 10
0.53
26.5
27
−45
A3
90 79 68 46.5
10 5 5 5
0.64
25.6
21.8
−32
B1
76 61 42 23
5 5 5 5
0.7
34
59
+29
B2
76 42 23 14
5 5 5 5
0.8
39.5
274
−19
B3
76 70 61 42
5 5 5 5
0.6
29.5
48
+53
1.3
Resistivity (m cm)
Stress (MPa)
Condition B1: strain gradient = 3.6 × 10−5 m−1 .
start to grow, but only a few reach the top. This results in a wide variation in grain size over the film thickness leading to a stress gradient through the layer. For the faster growing PECVD films [13] more columnar-like grains and thus a reduced stress gradi-
ent was observed compared to CVD SiGe for films with equal thickness. Based on this, we believe that the very high hydrogen dilution should be used in the beginning of the growth process, where nucleation has to be achieved and V-shaped grains start to grow. Afterwards, a reduced H2 dilution can be used for ‘tuning’ the shape of the grains, restricting them to grow more columnar-
Fig. 6. Cross-section TEM of 0.84 m thick cSiGe deposited using variable H2 dilution, Table 4, condition A1. Elongated columnar grains are observed, smaller grains at the SiGe-oxide interface, partially amorphous fraction.
Fig. 7. Cross-section TEM of 0.74 m thick cSiGe deposited using variable H2 dilution, Table 4, condition A3 (top of the picture: substrate side; bottom: film surface). Elongated columnar grains are observed.
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like and therefore giving us a possibility to tune the strain gradient. For condition B1 (Table 4) a very low strain gradient of 3.6 × 10−5 m−1 was measured, which is a factor of 10 better than the best value obtained using a compressive top layer at 350 ◦ C. 3.2. Multi-layer combining CVD and PECVD 3.2.1. Film characterization The microcrystalline deposition method results in excellent films at very low temperatures. Due to the hydrogen dilution however, a rather limited growth rate is obtained. Even in the case of the variable hydrogen dilution, the growth rate stays below 40 nm/min. Sometimes, e.g. for inertial sensors, the deposition of very thick layers (up to 10 m) is needed. The use of PECVD without hydrogen dilution would in that case be more economical. However, PECVD depositions alone at CMOS-compatible temperatures (∼450 ◦ C or lower) give amorphous films characterized by a high compressive stress and high resistivity values [3]. An LPCVD process gives crystalline films at these low temperatures, but has a very low deposition rate. Hence, it was shown in [3] that combining the two processes gives the benefits of both PECVD and CVD (without plasma), namely a high deposition rate and a low crystallization temperature, respectively. The polycrystalline SiGe multi-layer stack used in this work is composed of a thin PECVD seed layer (a-Si or a-SiGe), a LPCVD SiGe crystallization layer and a bulk PECVD SiGe layer (Fig. 8). The process time for the last PECVD step determines the
Fig. 8. Cross-section TEM for 1 m multi-layer film showing columnar grains (Si35 Ge65 ).
Fig. 9. Deflection of a 1 mm long released cantilever for 10 m thick SiGe layer with 800 nm Si-rich top layer.
thickness of the layer and the average deposition rate. While this process was introduced in a previous paper [3], in this work the strain gradient is further fine-tuned (down to 3.5 × 10−6 m−1 ) using a top compressive layer. A 1 m film exhibits a low average stress of −5 MPa, and an average resistivity of 1 m cm. It has a polycrystalline structure with columnar grains. Ten micrometers thick layers exhibit high deposition rates (∼100 nm/min), a low resistivity of 0.9 m cm, a tensile stress of 71 MPa and a positive strain gradient of 3.6 × 10−5 m−1 . 3.2.2. Strain gradient optimization Similar as shown for the microcrystalline films, it is possible to modify the stress and strain gradient in these multi-layers by stacking layers of opposing stress behavior. Thus, a top compensation layer on the multi-layer stack could further reduce the strain gradient. A few options that were explored for a 10 m thick film with a-SiGe seed layer are as follows: (a) a CVD top layer, (b) a top layer deposited at high plasma power and (c) a silicon rich top layer. Of these, a 10 m film with 1 m silicon-rich top layer gives the best results with a tensile stress of 57 MPa and a strain gradient of 1.2 × 10−5 m−1 [3]. The thickness and composition of the top compensation layer depends on the total
Fig. 10. Stress profile of a 11.6 m thick layer rebuilt from stress gradient measurements done after successive thinning down the layers using low energy plasma.
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Fig. 11. Thick poly-SiGe gyroscope on top of standard 0.35 m Al-CMOS process.
film stack built up (including the seed layer and the underlying sacrificial layer) and the stress gradient aimed for, as shown below. For example, a 10 m thick film with an a-Si seed layer and a 800 nm thick Si-rich compensation layer exhibited an average stress of 35 MPa, a resistivity of 1.45 m cm, a deposition rate of 90 nm/min and an extremely low strain gradient value (3.6 × 10−6 m−1 ) (Fig. 9). The amorphous seed layer is highly compressive and is expected to have a big influence on the overall strain gradient. This is apparent from the fact that a much thicker (1 m) top compensation layer is required to achieve a strain gradient of 1.2 × 10−5 m−1 for an a-SiGe seed layer (see above) as compared to an a-Si seed layer, where a 800 nm thick compensation layer gives a strain gradient of 3.6 × 10−6 m−1 . For a similar 10 m thick poly-SiGe layer, the stress profile across its thickness was derived. This was achieved by successively thinning down the layer, using very low energy plasma to avoid damaging the film, and measuring the strain gradient at each thickness. The strain gradient data was then used to construct the complete stress profile [14]. Fig. 10 shows that the maximum stress variation is in the initial layers after which the stress becomes more or less constant. The top compensation layer reverses the stress profile. Thus, at first it was expected that, while keeping the seed layer and top compensation layer constant, and just by varying the thickness of the bulk PECVD layer, it should be possible to obtain similar strain gradient values for different film thickness, with very little optimization required. Starting with a similar stack as for the 10 m layer, 4 m thick poly-SiGe films were developed. In this case however, the compensation layer needed to be thicker than expected to achieve
a low strain gradient. For a 1.35 m thick top layer, a stress value of 20 MPa, a resistivity of 1 m cm and a strain gradient of 3.5 × 10−6 m−1 is obtained. This result shows that, although the top compensation layer needed fine-tuning, it is possible to realize low strain gradient layers at different thickness. The optimized films described here can be used as structural layers for CMOS-integrated gyroscopes (Fig. 11) [15], accelerometers, etc. As these multi-layer films with top compressive layer have a variable Ge content over their layer thickness, the stability of these layers with time and temperature initially was a concern. In [6] it is however shown that both 4 and 10 m thick films are very stable over time and during temperature variations between room temperature and 100 ◦ C. 4. Conclusions The work presented above gives ample evidence that both microcrystalline SiGe and CVD/PECVD multi-layer poly-SiGe films form excellent MEMS structural layers, especially of interest for low thermal budget applications. Excellent electrical and mechanical properties were obtained with extremely low strain gradients (<2 m deflection for 1 mm long cantilever). These low strain gradients can be either reached through the use of top compressive layers for both types of films or through the use of a variable hydrogen dilution in the case of microcrystalline SiGe. These films are also stable with respect to time and temperature. First CMOS-integrated moving gyroscopes have been realized and other applications are currently being evaluated.
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Acknowledgements The work on 10 m thick films presented here is partially supported by the European project “SiGeM” (IST-2001-37681). The authors would also like to acknowledge the SiGeM project partners and Tim Brosnihan for the valuable discussions and suggestions. References [1] A. Witvrouw, M. Gromova, A. Mehta, S. Sedky, P. De Moor, K. Baert, C. Van Hoof, Poly-SiGe, a superb material for MEMS, Proc. MRS 782 (2004) 25–36. [2] R.T. Howe, T.J. King, Low-temperature LPCVD MEMS technologies, MRS Spring Symp. U 729 (2002) 205–213. [3] A. Mehta, M. Gromova, C. Rusu, O. Richard, K. Baert, C. Van Hoof, A. Witvrouw, Novel high growth rate processes for depositing poly-SiGe structural layers at CMOS compatible temperatures, in: Proceedings of the MEMS’04, 2004, pp. 721–724. [4] M. Gromova, K. Baert, C. Van Hoof, A. Mehta, A. Witvrouw, The novel use of low temperature hydrogenated microcrystalline silicon germanium for MEMS applications, Microelectron. Eng. 76 (2004) 266– 271. [5] T.J. King, R.T. Howe, S. Sedky, G. Liu, B.C. Lin, M. Wasilik, C. Duenn, Recent progress in modularly integrated MEMS technologies, Digest IEEE IEDM (2002) 199–202. [6] A. Mehta, M. Gromova, P. Czarnecki, K. Baert, A. Witvrouw, Optimization of PECVD poly-SiGe layers for MEMS post-processing on top of CMOS, in: Proceedings of the Transducers’05, 2005, pp. 1326– 1329. [7] S. Sedky, A. Witvrouw, A. Saerens, P. Van Houtte, J. Portmans, K. Baert, Effect of in situ boron doping on properties of SiGe films deposited by chemical vapor deposition at 400 ◦ C, J. Mater. Res. 16 (9) (2001) 2607–2612. [8] H.C. Lin, C.Y. Chang, W.H. Chen, W.C. Tsai, T.C. Chang, T.G. Jung, H.Y. Lin, Effects of SiH4 , GeH4 and B2 H6 on the nucleation and deposition of polycrystalline Si1−x Gex films, J. Electrochem. Soc. 141 (9) (1994) 2559–2563. [9] G. Ganguly, T. Ikeda, T. Nishimiya, K. Saitoh, M. Kondo, A. Matsuda, Hydrogenated microcrystalline silicon germanium: a bottom cell material for amourphous silicon-based tandem solar cells, Appl. Phys. Lett. 69 (27) (1996) 4224–4226. [10] J. Yang, H. Kahn, A.-Q. He, S.M. Phillips, A.H. Heuer, A new technique for producing large-area as-deposited zero-stress LPCVD polysilicon films: the multipoly process, J. Microelectromech. Syst. 9 (4) (2000) 485–494. [11] K. Tanaka, A. Matsuda, Glow Discharge Amorphous Si Growth Process and Structure, NHPC, Amsterdam, 1987, pp. 156–158. [12] K. Baert, Very Low Temperature Growth of Microcrystalline and Epitaxial Silicon by PECVD, KULeuven, Leuven, 1990, pp. 9– 10.
[13] T. Van der Donck, J. Proost, C. Rusu, K. Baert, C. Van Hoof, J.-P. Celis, A. Witvrouw, Effect of deposition parameters on the stress gradient of CVD and PECVD poly-SiGe for MEMS applications, Proc. SPIE 5342 (2004) 8–18. [14] A. Molfese, A. Mehta, A. Witvrouw, Determination of stress profile and optimization of stress gradient in PECVD poly-SiGe films, Sens. Actuators: Phys. 118 (2) (2005) 313–321. [15] A. Witvrouw, A. Mehta, A. Verbist, B. Du Bois, S. Van Aerde, J. RamosMartos, J. Ceballos, A. Ragel, J.M. Mora, M.A. Lagos, A. Arias, J.M. Hinojosa, J. Spengler, C. Leinenbach, T. Fuchs, S. Kronmuller, Processing of MEMS gyroscopes on top of CMOS Ics, Proc. ISSCC (2005) 88–89.
Biographies Maria Gromova received an MS degree in Communication Engineering from the Technical University of Sofia, Bulgaria in 1999. She is currently pursuing a PhD degree in Electronics Engineering at IMEC, Leuven, with main research interests including post-CMOS integration of MEMS. In particular hydrogenated microcrystalline SiGe development for MEMS applications. Anshu Mehta received her Bachelor’s degree from Indian Institute of Technology, Bombay in 2000, and a MS in engineering science from Pennsylvania State University in 2002. She has been working at Interuniversity MicroElectronics Center, Leuven, as a Research Engineer, since 2002. Her research interests include poly-SiGe surface micromachining, focusing on CMOSintegrated MEMS. Kris Baert received his PhD degree in microelectronics and materials science at Leuven University, Belgium in 1990. His dissertation was on lowtemperature plasma-enhanced CVD of amorphous and microcrystalline Si for solar cells. In 1988, he was a Visiting Scientist at the Konagai-Takahashi Laboratory, Tokio Institute of Technology, Japan. From 1990 to 1992, he was a Research Scientist at the Materials and Electronics Devices Laboratory, Mitsubishi Electric, Amagasaki, Japan, developing poly-Si TFTs. In 1992, he joined the Interuniversity Micro-electronics Center, Leuven, as a Senior Scientist working on micromachined IR sensors. Since 1995, he has been responsible for development of microsystem technologies. He is currently head of the MEMS group working on RF-MEMS, poly-SiGe surface micromachining and CMOS integration. Ann Witvrouw received the MS degree in metallurgical engineering in 1986 from the Katholieke Universiteit, Leuven, Belgium, and both the MS degree in applied physics in 1987 and PhD degree in applied physics in 1992 from Harvard University, Cambridge, MA. In 1992, she joined the Interuniversity Microelectronics Center (IMEC), Leuven, where she worked on the reliability of metal interconnects until the end of 1998. During this time, her research was focused on the mechanical stress in films and lines, electromigration and stress induced voiding. In 1998, she switched to research in Micro-electromechanical Systems at IMEC, where she is now responsible for advanced MEMS process technologies including the integration of MEMS and CMOS. She has been the coordinator of the IST project SUMICAP and currently she is the coordinator of the IST project SiGeM.