Background

High-quality surface passivation is very important for a range of crystalline silicon (c-Si)-based electronic devices, and especially for high-efficiency c-Si solar cells. As the need for lower-cost silicon solar cells increases, since Si material has a rather high cost, thinner Si substrates are required. Therefore, their surface/volume ratio of such substrates and the contribution of their surfaces to the overall performance are increasing. Traditional surface passivation for Si involves the formation of a thin silicon dioxide (SiO2) layer. However, this process requires a high thermal budget process, which involves long period at high temperature. Owing to these process-related issues, considerable efforts have been made in the past to develop low-temperature surface passivation methods for both heavily doped and moderately doped c-Si surfaces. Besides SiO2, other layers such as SiC, a-Si:H and Si3N4 have been used for surface passivation [1]. Recently, Al2O3 films that are grown by atomic layer deposition (ALD) have been demonstrated to provide good surface passivation on c-Si [2,3,4]. ALD technique is a powerful method. It provides a high-level degree of precise control over the properties of the material, and especially the morphology and thickness of dielectric layers.

In the advanced semiconductor industry, hafnium dioxide (HfO2) thin films are used to replace SiO2 as the gate dielectric in field-effect transistors because they have better functionality and performance at lower cost [5, 6]. Additionally, the high refractive index of HfO2 makes it a potential candidate for anti-reflection coatings [7] and interference filters [8]. However, its surface passivation properties, particularly on c-Si, have scantly been studied. For example, Jun Wang et al. [9] presented the surface passivation properties of a Si surface using a thin HfO2 layer grown by ALD without further annealing. In another study Huijuan Geng et al. [10] reported advanced passivation using simple materials (Al2O3, HfO2) and their compounds H(Hf)A(Al)O deposited by ALD. All of the previous attempts were performed to deposit HfO2 on c-Si substrates without any pre-treatments.

In this work, the surface passivation properties of the HfO2 films deposited by a remote plasma atomic layer deposition system (RP-ALD) on p-type c-Si with and without in situ oxygen plasma pretreatment were investigated. Samples were annealed at different temperatures by rapid thermal annealing (RTA) system. The structural changes and the electrical properties of the thin films induced by RTA were characterized by field-emission transmission electron microscope (FE-TEM), X-ray photoelectron spectroscopy (XPS) and capacitance-voltage (C-V) measurements. The passivation mechanism of HfO2 films on Si is also investigated.

Methods

In this study, (100) oriented boron-doped p-type crystalline Czochralski (Cz) Si wafers that were polished on both sides and had a resistivity of 30 Ω · cm, original lifetime of 1 μs and a thickness of 250 μm were used. Prior to the deposition of the HfO2 film, all wafers were cleaned through a standard Radio Corporation of America (RCA) cleaning process followed by a dip in diluted hydrofluoric acid (HF) solution (5%) for 2 min to remove the native oxide and dried in nitrogen.

The HfO2 thin films were grown in an RP-ALD reactor (Model: Picosun, Finland) using tetrakis (ethylmethylamino) hafnium (TEMAH) and remote O2 plasma as the precursors for hafnium and oxygen respectively with N2 as the carrier gas. In the ALD process, one deposition cycle consisted of two half cycles, one TEMAH pulse (for 1.6 s) and one O2 plasma pulse (for 10 s). The nitrogen purge times for TEMAH and O2 were 10.0 and 12.0 s, respectively. The samples were divided into two groups. For group one, HfO2 thin films were deposited directly on the cleaned Si wafers. For group two, before deposition of HfO2 thin films, Si wafers were additionally treated by remote O2 plasma for 1 min. The O2 plasma power for the pretreatment and for the ALD deposition process was 2500 W. The HfO2 films for all of the samples were deposited at 250 °C. Different HfO2 thickness (5, 15, and 25 nm) were prepared on as-cleaned Si wafers followed by annealing at 500 °C, and the corresponding minority carrier lifetimes of the passivated wafer were 9.98, 66.8, and 4.2 μs, respectively, at the injection level of 3 × 1014 cm−3. Therefore, the thickness of 15 nm (corresponding to 168 ALD cycles) was used. The substrate pre-treatment could affect nucleation, leading to different film thickness. The thicknesses of the deposited HfO2 are 15 nm ± 0.5 nm and 13 nm ± 0.7 nm for the samples with and without the oxygen plasma pretreatment, respectively. The wafer was flat on a platen. The double side coated samples processed twice, with a break in vacuum to flip the wafer in the chamber. The HfO2 thin films were deposited on 2-in wafers. As the substrate holder was about 8 in, four samples were placed on the holder and processed at a time. The samples in the two groups are referred hereafter as SD (direct depositing samples) and SO (O2 plasma pretreatment samples), respectively. Annealing process was performed using a RTA system at 400–650 °C in N2 ambient for 10 min. Samples were identified with suffixes A400 to A650 that represent the annealing temperatures. Table 1 lists the samples.

Table 1 Details of the HfO2 thin films
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The minority carrier lifetimes (τ eff) of the samples were assessed by photo-conductance decay method (Model: WCT-120, Sinton lifetime tester) in the quasi-steady state mode. Metal-insulator-semiconductor (MIS) structures were prepared by depositing Al electrodes with diameters of 500 μm onto the passivation layer using a sputter system and a shadow mask. The C-V characteristics were measured with a HP4284A semiconductor characterization system to extract the electrical parameters. The chemical composition and states of elements in the HfO2/Si were analyzed by XPS (Thermo Fisher K-Alpha). The ion energy used for the depth profile was 3000 eV. The physical thicknesses, microstructure and interface properties of the HfO2 thin films were determined by FE-TEM (JEM-2100 F).

Results and Discussion

Generally, the quality of passivation is assessed in terms of τ eff or surface recombination velocity (SRV = S max ). The τ eff refers to the recombination at surface defects. Figure 1(a) plots τ eff and S max for all samples at the injection level of 3 × 1014 cm−3. The τ eff measurements were performed three times for each sample in the different locations, and the errors of the minority carrier lifetime were within ±5%. As the annealing temperature was increased from 400 to 500 °C, the τ eff of the annealed HfO2 sample with O2 plasma pretreatment at the injection level of 3 × 1014 cm−3 increased significantly. The increase of the annealed HfO2 samples without O2 plasma pretreatment at the same injection level was much less than that of the annealed HfO2 samples with O2 plasma pretreatment. At lower temperatures (T < 500 °C), the annealed HfO2 samples without O2 plasma pretreatment had lower τ eff than those with O2 plasma pretreatment. The annealing process provides energy to the HfO2 layer to active the passivation. When the annealing temperature higher than 500 °C, the minority carrier lifetime decreases, which might be due to the defects generated by the increased microcrystalline fraction and grain boundaries in the HfO2 layer. The O2 plasma pretreatment sample that had been annealed at the temperature of 500 °C had the highest τ eff of 67 μs, corresponding to an S max value of 187 cm/s. This calculation was based on the quasi steady-state photo conductance (QSSPC) τ eff data for the injection level of 3 × 1014 cm−3. S max represents the upper limit of SRV, and is estimated from the measured lifetime values using the following relation [11],

Fig. 1
figure 1

a τ eff and S max of the samples at the injection level of 3 × 1014 cm−3. Injection level-dependent effective minority carrier lifetime of the (b) SD and (c) SO samples

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$$ {S}_{max}=\frac{W}{2{\tau}_{eff}}, $$
(1)

where W (=250 μm) is the thickness of the silicon substrate. The lower value of S max can be attributed to a lower density of interface traps. It also can be seen from Fig. 1a that the O2 plasma pretreatment samples exhibited better passivation than the directly deposited samples, so they had a lower interface recombination velocity. This difference is attributable to the diffusion of O from the O2 plasma to the interfacial region to form a SiO2 thin film, which provides better chemical passivation of the dangling bonds.

Figure 1b, c shows the injection level-dependent effective minority carrier lifetime of the samples without and with O2 plasma pretreatment. For the SO samples, the minority carrier lifetime increases with the annealing temperature between 400 and 500 °C. All of the SO samples without annealing exhibited almost no passivation, and their τ eff values were similar to that of the bare Si wafer. However, τ eff of the annealed samples increased significantly and then decreased as the injection level increased from 4 × 1013 cm−3 to 5 × 1015 cm−3. The drop in τ eff with increasing injection levels is caused by Auger recombination in the bulk of the c-Si substrate. The τ eff of the as-deposited samples depends very strongly on injection level, decreasing by approximately one order of magnitude as the injection level is decreased from 3 × 1014 to 1013 cm−3. This dependence in injection levels is much weaker for the annealed samples. The τ eff values of the annealed samples decrease only slightly as the injection level is reduced [12].

C-V measurements are commonly used to characterize the quality of dielectric layers and their interface with the substrates. C-V measurements were performed herein at room temperature in the dark conditions at 1 MHz on a standard MIS (Al/HfO2/p-Si) structure. Figure 2a, b shows the C-V curves of the HfO2 thin films without and with O2 plasma pretreatment, respectively. The voltage (V A ) that was applied across the MIS device was varied (-5 V < V A  < 5 V) with a sweep step length of 100 mV and signal amplitude of 50 mA, shifting from accumulation to inversion. The shift of C-V curves toward negative voltages demonstrates the presence of effective oxide charges of positive polarity in the as-deposited HfO2 thin films. The effective oxide charge represents the sum of mobile ionic charges (Q m ), oxide trapped charges (Q OT ) and oxide fixed charges (Q f ). Q f significantly affects the flat band voltage (V FB ), as it is located at the oxide-semiconductor interface. In Fig. 2a, the C-V curves are shifted in the positive direction by the V FB shift because Q f decreases as the annealing temperature increases. The slope of the C-V curve increases with the annealing temperature increases, indicating that the interface trap density decreases as the annealing temperature increases. The HfO2 thin films with O2 plasma pretreatment exhibited a similar relationship, as shown in Fig. 2b. The presence of fixed charges arose from the charged oxygen vacancies in the films [13]. The fixed charge density is estimated using Eq. (2), assuming a negligible effect of the interface traps [14],

Fig. 2
figure 2

C-V characteristics measured at 1 MHz for (a) directly deposited samples without O2 plasma pretreatment, and (b) samples with O2 plasma pretreatment; (c) estimated Q f of the annealed HfO2 thin films

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$$ {V}_{FB}={\phi}_{ms}-\frac{q{ Q}_f}{C_{ox}}, $$
(2)

where ϕ ms (=0.32 eV), q (=1.602 × 10−19 C), C ox , and V FB are the difference between the work functions of metal and the semiconductor, the electronic charge, the capacitance of the dielectric per unit area and the flat band voltage, respectively.

The values of Q f for the as-deposited and annealed HfO2 thin films are shown in Fig. 2c. Q f decreases as the annealing temperature increases. The annealing process appears to reduce the density of oxygen vacancies that are responsible for the presence of positive fixed charges, which may be related to the reconstruction of the oxide film near the interface [15]. Furthermore, the Q f of SO samples are lower than that of SD samples at the same annealing temperature. The interfacial defect density (D it ) is determined using an approximation method given by W. A. Hill and C. C. Coleman [16]. The Q f and D it values are listed in Table 2.

Table 2 Calculated fixed charge density (Q f ) and interface defect density (D it ) from C-V measurement of the HfO2 thin films
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Cross-sections of the annealed thin films were evaluated by a FE-TEM for assessing the film microstructure and HfO2/Si interface. The FE-TEM cross-section analysis of the HfO2 thin film annealed at 500 °C (a) without and (b) with O2 plasma pretreatment is shown in Fig. 3. From the FE-TEM images, the annealed HfO2 thin films consist of three regions, which are the HfO2 layer, an interfacial oxide and the Si substrate. The atoms in the HfO2 layer are orderly arranged in some areas, indicating that the HfO2 layer is microcrystalline structure. A very thin interfacial oxide layer is formed between the high k film and the substrate in the as deposited and annealed samples [17]. The HfO2 layer and the interfacial layer of the sample with oxygen plasma treatment are 15.3 and 2.7 nm, respectively. Whereas, the HfO2 layer and the interfacial layer of the sample without the pretreatment are 13.9 and 2.2 nm, respectively. This thickness difference should not cause the significant lifetime variation (35 and 67 μs for the samples without and with the pretreatment). Therefore, the significant lifetime improvement could be attributed to the different interface layers with the oxygen plasma pretreatment.

Fig. 3
figure 3

FE-TEM cross-section analysis of the HfO2 film annealed at 500 °C (a) without and (b) with O2 plasma pretreatment

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Figure 4 shows the elemental depth profiles of the HfO2 films annealed at 500°C without and with O2 plasma pretreatment obtained by XPS. Three regions are observed. In Region A, when the etching time was below 100 s, the relatively uniform atomic percentages of Hf and O corresponded to the RP-ALD μc-HfO2 layer. In Region B, the O and Hf atomic percentages decreased as the etching time increased from 130 to 175 s, indicating that the O elements diffused into the c-Si substrate, forming an interfacial layer [18, 19]. In Region C, when the etching time increased above 175 s, the Si signal drastically increased up to more than 60%, corresponding to the surface of the c-Si substrate. The oxygen atomic percentage and Hf atomic percentage in the c-Si substrates are due to the Ar ion sputtering effect. During the sputter process of the XPS measurement, some of the Hf or O atoms may reside on the silicon substrate surface and then be detected. Notice that in Region B, in addition to lower Hf and O with a corresponding increase in Si signal in the interface region, the sample with the oxygen pretreatment has also a larger Si signal in the bulk of the HfO2 film that may account for the percentage differences. The similar results can be obtained at the other investigated annealing temperatures. A possible reason might be that the O2 pretreatment leads to the growth of a very thin SiO2 layer reducing the Hf and O diffusion coming from the subsequently deposited HfO2. Fewer atomic vacancies are formed by diffusion in the HfO2 on the sample with the O2 pretreatment. Thus, the O2 pretreatment can be expected to yield fewer interface traps and exhibited higher chemical passivation quality.

Fig. 4
figure 4

Elemental depth profiles of HfO2 annealed at 500 °C without and with O2 plasma pretreatment versus etching time

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Growth of a thin oxide film on a clean but unpassivated Si surface leads to the formation of new covalent bonds (chemical passivation) and termination of the dangling bonds [9]. Si/oxide interfaces often carry some fixed charges. These charges can induce an electric field at the surface of Si and can potentially reduce the recombination rate at the Si/oxide interface (field effect passivation). It has been reported by Hoex et al [20] that when preparing Al2O3 thin films by plasma ALD, they found a very thin (~1.5 nm) SiOx interfacial oxide layer was formed, which provides good passivation to c-Si surface. They attributed this to the exposure of the substrate to the oxygen plasma in the very first ALD cycles. Although in this study the HfO2 thin films are prepared, the oxygen plasma pretreatment is found to result in a similar interfacial oxide layer (a-HfO2 + a-SiO2). The oxygen plasma pretreatment could improve the surface passivation of Si wafers.

Conclusions

In this work, HfO2 thin films with a thickness of 15 nm were deposited on p-type crystalline silicon wafers using a remote plasma atomic layer deposition system. In situ remote O2 plasma pretreatment of the Si substrate before the deposition of HfO2 thin films and post-annealing at 500 °C for 10 min effectively reduced the trap density at the HfO2/Si interface, yielding a highest lifetime of 67 μs. The HfO2 thin films deposited by RP-ALD with O2 plasma pretreatment have potential as passivation layers in high-quality Si solar cells.