**3. Channel mobility and interface state density**

Mobility is a measure of the ease a carrier can be moved in a solid under the application of an electric field *ξ*. It can be related to the speed of the carrier which is limited by scattering events occurring at average time intervals *τ*. Several types of scattering processes affect transport but the ones yielding the most frequent disturbances define the mobility value which can be written explicitly as

$$\frac{1}{\frac{1}{\mu}} = \frac{\frac{\pi}{\zeta}}{\bar{v}} = \frac{m^\*}{q} \sum\_{i} \frac{1}{\tau\_i} \tag{9}$$

10.5772/54396

257

http://dx.doi.org/10.5772/54396

atoms. Long Si-Si bonds, i.e. bigger than 2.35 Å, can have energy levels inside both the Si and SiC gaps. But the splitting between short Si-Si levels can be outside the silicon gap while being electrically active at a SiO2/SiC interface [90, 130], as illustrated in Fig. 4. The same analysis is true when comparing defect levels in the 6H- and 4H-SiC polytypes which have energy gaps of 3 and 3.3 eV, respectively. Because their valence bands are lined up, the conduction band of 4H-SiC is 0.3 eV higher. This is why the major defect affecting 4H-SiC, located 0.1 eV below its conduction band, has a limited impact on 6H-SiC devices [114]. It is interesting to note that this defect corresponds to an energy of 0.4 eV above the Si conduction band and that a similar trap level has been observed in SiO2 on Si even though it does not

The second reason explaining the large defect density at SiC interfaces is the oxidation

Like silicon oxidation, it follows the Deal-Grove reaction-diffusion model, so that the thermal

*B* 2 

where *B*/*A* is the linear rate constant, *B* is the parabolic rate constant, and *t*<sup>0</sup> is an offset time constant [45]. For details on SiC oxidation kinetics and parameters corresponding to various orientations, pressures, and temperatures, see Refs. [43, 68, 98, 113, 122, 127, 135, 136]. But unlike Si oxidation, Eq. (10) implies the release of carbon. Because of this complex multi-step process, a variety of atomic defects involving C can result from oxide formation if

Now that we have reviewed the impact and origin of defects at the SiO2/SiC interface, a comprehensive picture of associated trap levels limiting inversion mobility can be put forward, as shown in Fig. 4. Thermal oxidation of silicon carbide results in a SiC*x*O*<sup>y</sup>* inter-layer that includes threefold and fourfold coordinated Si and C atoms [8, 44, 90, 130]. Some generate dangling bonds whose energy spreads across the band gap because it is determined by the environment surrounding the defects. Another stable configuration yielding a trap level in the gap is a split C interstitial. When substituting for a Si site in the semiconductor, it can be viewed as a small C-aggregate or "C-cluster" comprising 6 atoms. A carbon-rich interface has indeed been observed by techniques such as Rutherford backscattering (RBS) [46, 57, 89, 97], x-ray photoelectron spectroscopy (XPS) [54, 62, 67, 77, 128], electron energy-loss spectroscopy (EELS) [33–35], Raman spectroscopy [80], and in-situ spectroscopic ellipsometry [63]. The dominant defects however, are likely oxide-related. Indeed, Si-Si bonds of various lengths can extend into SiO2 yielding interface and border traps [6, 17, 50, 72, 92, 112]. As mentioned before, the majority of corresponding bonding and antibonding states are located outside of the silicon gap but the former induce a distribution centered between the 4H and 6H-SiC conduction bands. The slightly smaller density of levels at lower energies is probably due to the majority of antibonding levels residing within the valence band of SiC whose edge is common in different polytypes [1]. The combination of trap levels associated with the SiC*x*O*<sup>y</sup>* inter-layer and Si-Si bonds, yields

1 (*B*/*A*)<sup>2</sup> <sup>+</sup>

2 *SiC* + 3*O*<sup>2</sup> → 2 *SiO*<sup>2</sup> + 2 *CO* (10)

Tailoring Oxide/Silicon Carbide Interfaces: NO Annealing and Beyond

4(*t* + *t*0)

*<sup>B</sup>* (11)

limit channel mobility [3, 4, 8].

oxide thickness as a function of time can be written as

*<sup>x</sup>*[*t*] = <sup>−</sup><sup>1</sup>

2 *B <sup>B</sup>*/*<sup>A</sup>* <sup>+</sup>

CO molecules do not all find their way to the gas phase. [48, 71, 90]

process

where *<sup>m</sup>*<sup>∗</sup> is the effective mass, *<sup>v</sup>*¯ is the net drift velocity, and *<sup>τ</sup><sup>i</sup>* corresponds to mean scattering times associated with various processes.

Because of the sudden termination of the semiconductor periodic lattice at the oxide interface, the channel mobility, *µch*, is expected to be lower than the bulk mobility. Indeed, electrically active levels can appear in the band gap and act as recombination centers or Coulomb scattering centers. Moreover, the free carriers can interact remotely with charged border states in the oxide, further reducing *τCb*. Other major damping mechanisms include surface-phonon and surface-roughness scattering represented by *τph* and *τsr*, respectively. We note that the coupling of carriers to scattering effects depends strongly on their velocity such that Coulomb scattering dominates at low fields, while surface roughness scattering becomes dominant at higher fields. Because of these interface phenomena, the best SiO2/Si devices display a channel mobility of about 700 cm2/V.s, equivalent to 50% of the bulk mobility [78]. In the case of 4H-SiC however, the native oxide interface yields mobilities of about 5 cm2/V.s, or less than 1% of the bulk value. So why is SiC so affected by the formation of a thermally grown interface? Let us discuss it from the point of view of Coulomb scattering, so that the question becomes: why is the density of interface traps (*Dit*) so much more prominent in SiC?

First, a down side of having a wider band gap is that it is sensitive to a wider range of defects. To first order, only the corresponding levels falling inside the band gap can be charged and yield Coulomb scattering. Since there is no evidence that SiO2 formed on SiC is any different from thermally oxidized silicon [99], it can contain the same type of defects, some having energy levels affecting only SiC carriers. Let's take the example of the oxygen vacancy, detected by electron spin resonance (ESR) in both systems, which yields Si-Si bonds [75, 106]. The energy split of that dimer is inversely proportional to the proximity between atoms. Long Si-Si bonds, i.e. bigger than 2.35 Å, can have energy levels inside both the Si and SiC gaps. But the splitting between short Si-Si levels can be outside the silicon gap while being electrically active at a SiO2/SiC interface [90, 130], as illustrated in Fig. 4. The same analysis is true when comparing defect levels in the 6H- and 4H-SiC polytypes which have energy gaps of 3 and 3.3 eV, respectively. Because their valence bands are lined up, the conduction band of 4H-SiC is 0.3 eV higher. This is why the major defect affecting 4H-SiC, located 0.1 eV below its conduction band, has a limited impact on 6H-SiC devices [114]. It is interesting to note that this defect corresponds to an energy of 0.4 eV above the Si conduction band and that a similar trap level has been observed in SiO2 on Si even though it does not limit channel mobility [3, 4, 8].

6 Physics and Technology of Silicon Carbide Devices

which can be written explicitly as

SiC?

**3. Channel mobility and interface state density**

scattering times associated with various processes.

1 *<sup>µ</sup>* <sup>=</sup> *<sup>ξ</sup>*

From Eqs. (7) & (8), it can be seen that the smaller the designed blocking voltage, the smaller the width of the necessary drift region, and the larger the contribution from the channel resistance. In Si, that does not have a major impact in power devices because *µch* can be as high as 50% of the bulk value *µdr*. However, SiC channels suffer from a mobility that can be less than 1% of *µdr* at the native SiO2/4H-SiC interface. Therefore, interface quality can affect performance even in the kV range and the full potential of the SiC material cannot be reached. This is highlighted in Fig. 1, where the ratio of ON resistances was calculated using Fig. 2(a) constants, *<sup>L</sup>* <sup>=</sup> <sup>1</sup>*µm*, *<sup>P</sup>* <sup>=</sup> <sup>10</sup>*µm*, *<sup>n</sup>* <sup>=</sup> 1015 cm<sup>−</sup>2, and the following SiC *<sup>µ</sup>ch*: 5, 50, and 500 cm2/V.s. The significance of those mobility values are discussed in the next Sections.

Mobility is a measure of the ease a carrier can be moved in a solid under the application of an electric field *ξ*. It can be related to the speed of the carrier which is limited by scattering events occurring at average time intervals *τ*. Several types of scattering processes affect transport but the ones yielding the most frequent disturbances define the mobility value

*<sup>v</sup>*¯ <sup>=</sup> *<sup>m</sup>*<sup>∗</sup>

where *<sup>m</sup>*<sup>∗</sup> is the effective mass, *<sup>v</sup>*¯ is the net drift velocity, and *<sup>τ</sup><sup>i</sup>* corresponds to mean

Because of the sudden termination of the semiconductor periodic lattice at the oxide interface, the channel mobility, *µch*, is expected to be lower than the bulk mobility. Indeed, electrically active levels can appear in the band gap and act as recombination centers or Coulomb scattering centers. Moreover, the free carriers can interact remotely with charged border states in the oxide, further reducing *τCb*. Other major damping mechanisms include surface-phonon and surface-roughness scattering represented by *τph* and *τsr*, respectively. We note that the coupling of carriers to scattering effects depends strongly on their velocity such that Coulomb scattering dominates at low fields, while surface roughness scattering becomes dominant at higher fields. Because of these interface phenomena, the best SiO2/Si devices display a channel mobility of about 700 cm2/V.s, equivalent to 50% of the bulk mobility [78]. In the case of 4H-SiC however, the native oxide interface yields mobilities of about 5 cm2/V.s, or less than 1% of the bulk value. So why is SiC so affected by the formation of a thermally grown interface? Let us discuss it from the point of view of Coulomb scattering, so that the question becomes: why is the density of interface traps (*Dit*) so much more prominent in

First, a down side of having a wider band gap is that it is sensitive to a wider range of defects. To first order, only the corresponding levels falling inside the band gap can be charged and yield Coulomb scattering. Since there is no evidence that SiO2 formed on SiC is any different from thermally oxidized silicon [99], it can contain the same type of defects, some having energy levels affecting only SiC carriers. Let's take the example of the oxygen vacancy, detected by electron spin resonance (ESR) in both systems, which yields Si-Si bonds [75, 106]. The energy split of that dimer is inversely proportional to the proximity between

*<sup>q</sup>* ∑ *i*

1 *τi*

(9)

The second reason explaining the large defect density at SiC interfaces is the oxidation process

$$2\text{SiC} + 3\text{O}\_2 \rightarrow 2\text{SiO}\_2 + 2\text{CO} \tag{10}$$

Like silicon oxidation, it follows the Deal-Grove reaction-diffusion model, so that the thermal oxide thickness as a function of time can be written as

$$\ln[t] = -\frac{1}{2}\frac{B}{B/A} + \frac{B}{2}\sqrt{\frac{1}{(B/A)^2} + \frac{4(t+t\_0)}{B}}\tag{11}$$

where *B*/*A* is the linear rate constant, *B* is the parabolic rate constant, and *t*<sup>0</sup> is an offset time constant [45]. For details on SiC oxidation kinetics and parameters corresponding to various orientations, pressures, and temperatures, see Refs. [43, 68, 98, 113, 122, 127, 135, 136]. But unlike Si oxidation, Eq. (10) implies the release of carbon. Because of this complex multi-step process, a variety of atomic defects involving C can result from oxide formation if CO molecules do not all find their way to the gas phase. [48, 71, 90]

Now that we have reviewed the impact and origin of defects at the SiO2/SiC interface, a comprehensive picture of associated trap levels limiting inversion mobility can be put forward, as shown in Fig. 4. Thermal oxidation of silicon carbide results in a SiC*x*O*<sup>y</sup>* inter-layer that includes threefold and fourfold coordinated Si and C atoms [8, 44, 90, 130]. Some generate dangling bonds whose energy spreads across the band gap because it is determined by the environment surrounding the defects. Another stable configuration yielding a trap level in the gap is a split C interstitial. When substituting for a Si site in the semiconductor, it can be viewed as a small C-aggregate or "C-cluster" comprising 6 atoms. A carbon-rich interface has indeed been observed by techniques such as Rutherford backscattering (RBS) [46, 57, 89, 97], x-ray photoelectron spectroscopy (XPS) [54, 62, 67, 77, 128], electron energy-loss spectroscopy (EELS) [33–35], Raman spectroscopy [80], and in-situ spectroscopic ellipsometry [63]. The dominant defects however, are likely oxide-related. Indeed, Si-Si bonds of various lengths can extend into SiO2 yielding interface and border traps [6, 17, 50, 72, 92, 112]. As mentioned before, the majority of corresponding bonding and antibonding states are located outside of the silicon gap but the former induce a distribution centered between the 4H and 6H-SiC conduction bands. The slightly smaller density of levels at lower energies is probably due to the majority of antibonding levels residing within the valence band of SiC whose edge is common in different polytypes [1]. The combination of trap levels associated with the SiC*x*O*<sup>y</sup>* inter-layer and Si-Si bonds, yields the U-shaped D*it* distribution. It rises sharply towards the SiC band edges because of the Si-related defects not dominant at silicon interfaces. Therefore, the efficiency of passivation techniques are expected to be very different at interfaces formed on the two semiconductors. 10.5772/54396

259

http://dx.doi.org/10.5772/54396

The benefits of NO annealing have been directly correlated with the incorporation of nitrogen, which is confined to the SiO2/SiC interface, as detected by various techniques such as secondary ion mass spectroscopy (SIMS) [82, 106], nuclear reaction analysis (NRA) [81], electron energy loss spectroscopy (EELS) [33], and medium energy ion scattering (MEIS) [47, 137]. To study the impact of nitrogen, the amount incorporated can be tailored by the NO annealing time as illustrated in Fig. 5(a). The N density is then extracted by integrating SIMS interface peaks resulting from 1175 ◦C NO exposure of a dry thermal oxide for up to 2 hours. The nitrogen content is found to saturate around 6 × <sup>10</sup><sup>14</sup> cm−<sup>2</sup> or about a half monolayer coverage of the SiC surface. The nitridation kinetics result from a balance between N incorporation and removal. Indeed, at this temperature, NO decomposes partially into N2 and O2. While 1175 ◦C is required to enable NO diffusion to the interface and subsequent nitridation, the presence of oxygen limits its effect as interfacial nitrogen is unstable against the slow re-oxidation occurring in parallel [37]. Moreover, additional defect formation can

Tailoring Oxide/Silicon Carbide Interfaces: NO Annealing and Beyond

Progressive reduction of the *Dit* across the 4H-SiC band gap corresponding to the tailored introduction of nitrogen has been measured in metal-oxide-semiconductor capacitors (MOSCAPs), as shown in Fig. 5(b). The density of states shows a strong correlation to the nitrogen content and is reduced by up to an order of magnitude close to the conduction band edge [82, 106, 108]. The sensitivity of the inversion mobility of electrons to the *Dit* reduction was studied in lateral field-effect transistors containing different amounts of nitrogen at the SiO2/4H-SiC interface. From the results depicted in Figs. 6(a) & 10, the peak field-effect mobility is found to be inversely proportional to the density of charged states, which reveals a Coulomb-scattering-limited transport. It is important to note that this is true even in devices with the lowest *Dit* so that further defect passivation is projected to increase the mobility from 50 to more than 100 cm2/V.s, which cannot be achieved by NO POA alone as nitrogen density becomes saturated. These conclusions are in agreement with separate mobility studies using Hall effect measurements on nitrided samples [13, 114, 125]. Such experiments also reveal that at higher fields, mobility becomes limited by surface-roughness scattering. Although NO POA has been shown to yield smoother interfaces [51], it is not clear what else can be

C and O Intensity (a.u.)

**Figure 5.** (a) SIMS Nitrogen profiles showing progressive accumulation at the SiO2/SiC interface with increasing NO annealing time. Adapted from Ref. [106]. (b) Density of interface traps across the 4H-SiC band gap. Longer NO anneals yield lower D*it*.

10<sup>11</sup>

Conduction Band

0.5 1.0 1.5 2.0 2.5 3.0 Ec - E (eV)

 As-Oxidized 7.5 min NO 30 min NO 4 hr NO

Valence Band

10<sup>12</sup>

Dit (cm -2. eV -1)

10<sup>13</sup>

Carbon

also result from the presence of the excess oxygen.

done to further reduce that particular component.

 **Nitrogen**

< < < SiO2 | SiC > > >

200 300 400 500 600 700 Depth (A)

Reproduced with permission from Ref. [108].©2011 IEEE

**(a) (b)**

7.5 min NO

30 min NO

2 hr NO

1.5

1.0

Oxygen

Nitrogen Concentration (1021cm-3

)

0.5

0.0

In the following Sections, we will discuss how to reduce *Dit* and its relationship with mobility. Although there is extensive literature dedicated to various orientations and polytypes, this overview is dedicated to devices fabricated on the (0001) Si-face of 4H-SiC.
