- Email: [email protected]
Effect of bias on the structure and properties of TiZrN thin films deposited by unbalanced magnetron sputtering
Effect of Bias on the Structure and Properties of TiZrN Thin Films Deposited by Unbalanced Magnetron Sputtering Yu-Wei Lin, Hsi-An Chen, Ge-Ping Yu, Jia-Hong Huang PII: DOI: Reference:
S0040-6090(16)30172-9 doi: 10.1016/j.tsf.2016.05.021 TSF 35204
To appear in:
Thin Solid Films
Received date: Revised date: Accepted date:
10 November 2015 16 May 2016 16 May 2016
Please cite this article as: Yu-Wei Lin, Hsi-An Chen, Ge-Ping Yu, Jia-Hong Huang, Effect of Bias on the Structure and Properties of TiZrN Thin Films Deposited by Unbalanced Magnetron Sputtering, Thin Solid Films (2016), doi: 10.1016/j.tsf.2016.05.021
This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Effect of Bias on the Structure and Properties of TiZrN Thin Films Deposited by
1
SC R
IP
Yu-Wei Lin1,3, Hsi-An Chen 1, Ge-Ping Yu1,2,*, Jia-Hong Huang1
T
Unbalanced Magnetron Sputtering
Department of Engineering and System Science, National Tsing Hua University, Taiwan,
R.O.C.
Institute of Nuclear Engineering and Science, National Tsing Hua University, Taiwan,
NU
2
3
MA
R.O.C.
Instrument Technology Research Center, National Applied Research Laboratories, Taiwan,
Presenter: Ge-Ping Yu
*
TE
†
D
R.O.C
AC
CE P
Corresponding author’s e-mail: [email protected]
ACCEPTED MANUSCRIPT ABSTRACT The objective is to investigate the substrate bias effect on the structure and properties
T
of the TiZrN thin films. The TiZrN thin films were deposited by direct current unbalanced
IP
magnetron sputtering system with dual guns (Ti and Zr) targets onto Si (100) substrates at
SC R
different substrate bias ranging from －35 V to －150 V. Experimental results indicated that all the specimens have strong (111) texture in X-ray Diffraction patterns. In this study, we
NU
discovered a transition bias of －35 V, above which a significant improvement of properties was found, including high hardness, excellent brilliance, low resistivity and fine surface
MA
morphology. Within the bias range of －40 to －120 V, 1the hardness of TiZrN films is around 35.5 GPa, the resistivity is about 33.5 μΩ-cm, and the brilliance is larger than 80. The
D
roughness is between 0.5 nm and 0.6 nm. The TiZrN films maintain excellent properties
TE
through a large range of applying bias, indicating that the process window is considerable wide. However, structure damage and thin film delamination were found when substrate bias
CE P
reached －150 V. The Rutherford backscatterng spectrometry result and scanning electron microscope image further support the structure damage at －150 V. For protective coatings,
AC
low residual stress is required to avoid delamination. By adjusting substrate bias, residual stress can be controlled to lower value. In this study, the residual stress of TiZrN films gradually decreases with decreasing the substrate bias ranging from －65 V to －35 V. The TiZrN thin films with high hardness, lower residual stress could be obtained simultaneously at low substrate bias of －40 V and －45 V, the hardness and residual stress are 33.4～34.5 GPa and －2.7 ～－3.7 GPa, respectively. Keywords: TiZrN, hardness, residual stress
ACCEPTED MANUSCRIPT Introduction Owing to their excellent mechanical properties, low resistance, golden color, and
T
superior corrosion resistance, titanium nitride (TiN) and zirconium nitride (ZrN) have been
IP
widely used in various industries over the past 20 years[1-4]. Their applications include
SC R
decorative coatings, diffusion barriers in microelectromechanical systems, and hard coatings on tools. However, the demand for harder and tougher materials driven by technological progress necessitates the development of ternary transition-metal nitride films such as TiZrN,
NU
TiSiN, and TiAlN. Compared with TiN and ZrN, ternary TiZrN films not only exhibit a low
MA
resistivity and golden color but also feature superior mechanical properties and better corrosion resistance[5, 6].
D
Previous study [7] indicated that the addition of Zr into TiN or Ti into ZrN could
TE
enhance the fracture toughness of hard coatings. This approach of incorporating a third element is effective for improving various properties of thin films. In general, the high
CE P
hardness of hard coatings usually accompanies high residual stress in films deposited via the physical vapor deposition (PVD) process. However, the presence of such high residual stress
AC
in thin films can result in reduced adhesion and film spallation. The residual stress can be reduced while retaining the excellent mechanical properties of films by controlling the substrate bias during deposition[8]. The substrate bias is an important parameter in the PVD process. It governs the total energy and momentum of adatoms while influencing the structure and properties of the resulting thin films. In addition, the substrate bias is easily adjusted during thin film deposition. Therefore, we chose to change the substrate bias as a single-variable experimental parameter. Our objective was to investigate the effect of substrate bias on the structure and properties of deposited TiZrN thin films. The influence of substrate bias on the microstructure, surface morphology, hardness, residual stress, resistivity, and packing factor
ACCEPTED MANUSCRIPT was investigated. Experiments
T
In this research, ternary TiZrN thin films were deposited using a reactive direct current
IP
unbalanced magnetron (UBM) sputtering sputtering system. p-type Si(100) wafers were
SC R
chosen as the substrate material; the dimensions of the Si substrates were 3.5 × 3.5 × 0.525 cm3. The sputtering sources were a Zr target (99.9%) and a Ti target (99.995%), each with a diameter of 2 inches.
NU
First, before the coating process, the specimens were ultrasonically cleaned sequentially in acetone and methanol for 5 min each. The chamber was evacuated to a base pressure of 6.7
MA
× 10−4 Pa using a pumping system at 400°C. The substrate was then pre-sputtered with Ar+
D
ions at a bias of −1000 V for 5 min with the Ar (99.9995% purity) flow rate fixed at 50 sccm
TE
to remove the surface oxide layer.
After pre-sputtering, the working gas (Ar) and reactive gas (N2) were introduced into the
CE P
chamber. The Ar gas flow rate was fixed at 30 sccm with a maximum gas flow rate of 50 sccm; the N2 gas flow rate was fixed at 1.9 sccm with a maximum gas flow rate of 5 sccm.
AC
The total gas pressure was maintained at 1.5 × 10−1 Pa. The coating temperature was 400°C and was monitored using a thermocouple positioned between the substrate holder and the halogen lamps. During deposition, the dual gun current was kept constant at 0.27 A. The sample number was assigned according to the substrate bias (e.g., sample B35 = 35 V substrate negative bias). The substrate bias was varied from −35 V to −150 V. The target-to-substrate distance was adjusted to 10 cm. The TiZrN thin films were deposited by co-sputtering of Ti and Zr targets. The specimens were rotated during sputtering. Both normal lines of the Ti and Zr target surfaces intersected on the revolved substrate surface to ensure chemical uniformity of the coating. The deposition time of all specimens was 60 min. The optimum parameters were determined by the Taguchi method[9].
ACCEPTED MANUSCRIPT All the deposited specimens were analyzed. The surface compositions of TiZrN films were determined with a PHI 1600 X-ray photoelectron spectroscope at a vacuum of 10-8 Pa.
T
The X-ray source was Mg Kα (1253.6 eV). In order to remove the surface oxide, prior to
IP
analysis, the specimen surface was pre-sputtered by 3 keVAr+ for 3 minutes. The
SC R
deconvolution and fitting of X-ray photoelectron spectroscopy (XPS) spectrum was executed by a computer program XPSPEAK 4.1. The background of the spectra was subtracted by Shirley method. The high resolution spectra were then deconvoluted and processed by
NU
least-square fitting procedure using Gaussian function. The atomic percentages of Ti, Zr, N
MA
and O were obtained from the integrated areas of deconvoluted spectra divided by the corresponding sensitivity factors. The relative content of each element was calculated by the 100%, where i represents Ti, Zr, N, or O, Ci is the content of
D
following equation: Ci=
TE
element i in the thin films, and I is the integrated intensity of the element i divided by the corresponding sensitivity factor. The X-ray source was Mg Kα (1253.6 eV). The N/(Ti+Zr)
CE P
ratio and packing factor were determined by Rutherford backscattering spectrometry (RBS). The Van de Graaff accelerator at NTHU was applied. The incident ion source was α (4He+)
AC
particles and the operating energy was 2 MeV. The chamber was evacuated below 8 × 10-4 Pa. A surface barrier detector was set up at the backscattering angle of 160° to the surface normal of specimens. Prior to analysis, the RBS was calibrated using Si and W standard samples. The resolution of the instrument was about 30 ~100 Å. The crystal structure and preferred orientation of TiZrN films were characterized by X-ray diffraction (XRD) using a D8 diffractometer. The Cu Kα line at 1.5405 Å was used as the source for diffraction. For the θ/2θ scans, a scan rate of 4deg/min was used, and the 2θ scan range was from 20 to 60
I ( hkl ) degrees. The texture coefficient of (hkl) orientation was defined as I ( hkl ) , where I is the integrated intensity of the corresponding (hkl) peak. The thickness and morphology of the
ACCEPTED MANUSCRIPT TiZrN thin films were observed by a JOEL JSM-6700F scanning electron microscopy (SEM) at National Chung Hsing University. The cross-sectional microstructure of TiZrN/Si was also
T
observed in SEM image. Before loading, the TiZrN/Si samples were cleaved into a piece of
IP
0.3×0.3 cm2 and attached on a sample holder with carbon tape and then the samples were
SC R
coated an Au layer to improve electric conductivity. The surface roughness and morphology of TiZrN thin films were observed by an atomic force microscope (AFM) with a silicon nitride tip. During the analysis, the contact mode which means that a constant contact force
NU
applied upon the specimen surface was chosen. A scanned frequency of 4 Hz line by line and
MA
2 × 2 μm2 scan size were set. The residual stress in the TiZrN thin films was measured using the laser curvature method. The electrical resistivity of the films was measured by the
D
four-point probe method. The coloration of the TiZrN coatings was characterized at room
TE
temperature using a Hunter Lab Miniscan XE Plus spectrophotometer, model 4000VSAV. The hardness of the deposited films was determined by nanoindentation tests
CE P
performed using a triangular Berkovich indenter. Sample hardness and elastic modulus were inferred from the results of the indentation tests. The tip was moved laterally across a surface
AC
while the normal force was varied or held constant. Displacement into the surface and lateral (tangential) force were measured during the process. The film hardness is theoretically a constant value but will vary with changes in the applied load. In general, the indentation depth should be less than one tenth of the film thickness to avoid the substrate effect. To evaluate the creep effect, the load was first increased over a period of 20 s to the desired load, maintained for 5 s, and then unloaded over a period of 20 s. The unloading stiffness was interpreted by the equation
S
dP 2 Er A dh
ACCEPTED MANUSCRIPT where the unloading stiffness S = dP/dh is measured experimentally, A is the projected area of the elastic contact, and Er is the reduced modulus.
IP
SC R
1 (1 2 ) (1 i2 ) Er E Ei
T
The reduced modulus is defined as
where E and Ei are the Young’s moduli for the specimen and indenter, respectively, and ν and νi are the Poisson’s ratios for the specimen and indenter, respectively. The reduced modulus
NU
can be determined if the contact area can be independently measured. Parameter A is a
MA
function of the contact depth (hc), and hc can be calculated by the following equation
hc hmax hs
D
where hmax is the total displacement determined from experimental data and hs is the depth of
TE
the surface at the surrounding of the contact:
( 2)
CE P
hs
(h h f )
2( 2) Pmax S
AC
where hf is the final depth of after load releasing and Pmax is the maximum load. Thus,
hc hmax 0.72
Pmax S
Finally, the hardness can also be determined as
H
Pmax A
where H is the film’s hardness, Pmax is the maximum load, and A is the projected area of contact. Before measurement of the coating samples, a quartz sample should be used to calibrate the Cx value of A:
A(hc ) C0 hc2 C1hc1 C2 hc1 / 2 C3 hc1 / 4 C4 hc1 / 6 C5 hc1 / 8
ACCEPTED MANUSCRIPT where C0 is a constant value of 24.5 for the Berkovich indenter, and C1–C5 values can be fitted using five different loads and contact depths. Notably, the hardness calibration value
T
and the reduced modulus of the quartz should be approximately 10 (±3%) GPa and 69.6
SC R
IP
(±3%) GPa, respectively.
Results and Discussion
According to XPS and RBS analyses, the N/(Ti+Zr) ratio for all samples is fixed at
NU
0.9–1.1, and the Zr/(Ti+Zr) ratios are approximately 0.55 (Table 1). The chemical
MA
composition of TiZrN thin films is not sensitive to substrate bias. An increase of the substrate bias not only increases the total energy but also enhances the momentum of the adatoms. The
D
momentum of an adatom is directly proportional to its mobility. Adatoms can move to stable
TE
sites as they obtain sufficient momentum, resulting in deposition of a denser thin film. By contrast, an excessively high negative substrate bias results in negative influences such as the
CE P
production of additional lattice defects, radiation damage, and thin film delamination. Fig. 1 presents SEM cross-sectional images showing the microstructure and thickness of
AC
the B40 specimens. The thickness of the TiZrN films ranges from 418 to 531 nm. The deposition rate was calculated as the thickness divided by the deposition time. The results demonstrate that the deposition rate decreases with increasing substrate bias because of the re-sputtering effect. The variation of deposition rate under different substrate biases is plotted in Fig. 2. Fig. 3 reveals that the packing factor of the films gradually increases from 0.82 to 0.96 when the substrate bias is between −35 V and −65 V; it then slightly decreases when the substrate bias exceeds −65 V. The aforementioned SEM cross-sectional results also correspond well with the packing factor. Fig. 3 also shows that the B35 and B150 specimens have lower packing factors.
ACCEPTED MANUSCRIPT Figs. 4 and 5 show the SEM surface images and AFM 3D images. The roughness of all samples is also given in Table 2. When the substrate bias exceeds −35 V, the adatoms obtain
T
sufficient momentum to move to stable sites; consequently, the TiZrN films become very
IP
smooth. Their roughness ranges from 0.5 nm to 0.8 nm. Although sample B150 also exhibits
SC R
a flat surface, a region of the B150 specimen delaminated because of high-energy ion bombardment, decreasing the adhesion strength (Fig. 6).
Fig. 7 shows XRD patterns of the TiZrN films deposited under different substrate biases.
NU
In XRD patterns, the reflection of the TiZrN (111) peak is observed between those of the TiN
MA
(111) peak and the ZrN (111) peak; similarly, the reflection of the TiZrN (200) peak is observed between those of the TiN (200) peak and the ZrN (200) peak. Moreover, according
D
to Hume–Rothery rules, TiZrN should be a substitutional solid solution. The difference
TE
between the diameters of the atoms (Zr and Ti) is less than 15%, the crystal structures of TiN and ZrN are both face-centered cubic, the electronegativities of Ti and Zn do not substantially
CE P
differ, and they both belong to the same group (IVB) in the periodic table. Thus, TiN and ZrN satisfy the generalized definition for a substitutional solid solution. The texture coefficient
AC
can be obtained by ITiZrN(111)/(ITiZrN(111) + ITiZrN(200)), where I(111) and I(200) represent the integrated intensities of the corresponding (111) and (200) peaks, respectively. Actually, all the specimens have strong (111) texture that is greater than or equal to 0.77. The B35 specimen has a lower (111) texture coefficient at −35 V, possibly because its looser structure results in random crystalline orientation. Therefore, the (111) texture coefficients increase with increasing substrate bias. Furthermore, Abadias proposed that texture variation can be attributed to a re-nucleation mechanism[10]. A high substrate bias results in the formation of more defects in the resulting thin films, and regions with a high defect density are prone to secondary nucleation and grain-boundary migration. Therefore, more grains grew along the (111) direction in the films grown in this study.
ACCEPTED MANUSCRIPT The full-width at half-maximum (FWHM) of TiZrN (111) as a function of substrate bias is shown in Fig. 8, where a smaller FWHM indicates good crystallinity of the TiZrN film. Fig.
T
9 illustrates variation of grain size as a function of substrate bias. The vertical dimension of
IP
the grains in TiZrN can be calculated using the Scherrer equation. The grain size of TiZrN
SC R
films for all samples is less than 22 nm. The FWHM(111) and grain size results are summarized in Table 2. The B35–B150 specimens were divided into three regions for purposes of discussion:
NU
Region 1: Low-negative-bias region with substrate bias ranging from −35 V to −65 V Region 2: Medium-negative-bias region with substrate bias ranging from −75 V to −95 V
MA
Region 3: High-negative-bias region with substrate bias ranging from −120 V to −150 V
Originally, the large grain size was approximately 20 nm in region 1. As the substrate
D
bias exceeded −65 V (entering region 2), the grain size decreased to approximately 8.3 nm.
TE
This decrease in grain size is attributed to the presence of additional defects induced by ion
CE P
bombardment. Areas with a high defect density are prone to nucleation; consequently, nucleation increases simultaneously with increasing number of defects, which decreases the grain size. The FWHM of (111) suddenly increased by 0.6° because the number of defects
AC
increased when the substrate bias exceeded −65 V. In region 3, the FWHM of (111) slightly decreased, possibly because the grains lacked sufficient energy for rearrangement of the (111) orientations at high substrate bias. Thus, the crystallinity of the films was enhanced. We observed similar behavior in our previous study[11]. According to the Scherrer equation, the grain size in the present study increased to approximately 14.5 nm from 8.3 nm. Compared to TiN and ZrN films, TiZrN films possess higher residual stress, which is related to the difference between the atomic radii: the atomic radius of Zr is larger than that of Ti. The formation of a substitutional solid solution by the addition of Zr into the TiN lattice or by the addition of Ti into the ZrN lattice induces serious lattice distortion, resulting in high residual stress. Fig. 10 shows the variation of the residual stress as a function of
ACCEPTED MANUSCRIPT substrate bias. It shows that the residual stress gradually increases from −0.8 to −5.5 GPa in region 1. The impact of high-energy Ar+ ions on a film surface during the deposition process
T
increases the residual stress in the film. The residual stress slightly decreased with increasing
IP
substrate bias in region 2, where the increase in the mobility of adatoms to stable sites
SC R
released the residual stress. When the substrate bias exceeded −95 V (region 3), the residual stress slightly increased because additional lattice defects formed under the high negative bias.
NU
Fig. 11 presents the variation of the electrical resistivity as a function of the substrate
MA
bias. The results show that all the samples exhibited low resistivity (from 30 to 52 μΩ·cm). Initially, the electrical resistivity suddenly decreased when the substrate bias exceeded −35 V in region 1. It thereafter remained constant (from 30 to 37 μΩ·cm) as the substrate bias was
TE
D
increased from −40 to −120 V. The electrical resistivity then slightly increased when the substrate bias exceeded −150 V.
CE P
The resistivity of the films is related to their number of lattice defects. An increase in the number of lattice defects can impede the movement of free electrons and increase their mean
AC
free path. Consequently, more lattice defects exist in the film, which may increase the film’s resistivity. The resistivity rapidly decreased when the substrate bias exceeded −35 V, which we attribute to the increase in densification. Increasing densification of the film results in a decrease in the number of lattice defects, further decreasing the resistivity of the thin films. However, in the case where the substrate bias was more negative than −95 V (region 3), the corresponding increase in resistivity is attributed to the fact that high-energy ion bombardment induced structural damage, resulting in a slight increase in the resistivity of the TiZrN films because of the introduction of additional defects. Therefore, we plotted the resistivity as a function of the packing factor in Fig. 12. The resulting plot indicates that the lowest packing corresponds to the highest resistivity.
ACCEPTED MANUSCRIPT Fig. 13 shows that the hardness of all specimens varied with the substrate bias. The hardness values of all the TiZrN films exceeded 33 GPa, ranging from 33.2 to 37.8 GPa,
T
except in the case of sample B35. Sample B35 exhibited the lowest hardness because of its
IP
looser structure, consistent with the SEM cross-section analysis and RBS results. The reason
SC R
the hardness of TiZrN films exceeds 33 GPa is likely associated with solid-solution strengthening. The addition of Zr into the TiN lattice or the addition of Ti into the ZrN lattice induces lattice distortion. The lattice distortion impedes the movement of dislocations and
NU
grain boundaries, thereby enhancing the hardness of TiZrN films.
MA
Moreover, texture, residual stress, grain size, and packing factor also influence the hardness of the TiZrN films. In NaCl-type TiZrN, the {110} planes and <110> direction
D
represent an active slip system. If the external force is parallel to the (111) plane normal, the
TE
Schmid factor is equal to zero on all slip systems[12]. Therefore, <111> is the hardest orientation, and the hardness of the film is enhanced with increasing (111) texture coefficient.
CE P
Similar conditions have been observed in the TiN system[13]. However, in our case, the hardness of the TiZrN films is not sensitive to changes in the (111) texture coefficient. The
AC
hardness of the TiZrN films might be affected by other factors such as residual stress, grain size, and packing factor. In this research, the grain size ranges from 8 nm to 21 nm for all samples. According to dislocation theory[14], the number of dislocations in a grain can be calculated; the authors of a previous study proposed a similar calculation for TiN thin films[15]. We can use the following equation to calculate the number of dislocations n in a grain:
where k is a factor that approaches unity, τ is the critical resolved shear stress, D is the grain size, G is the shear modulus, and b is the Burgers vector.
ACCEPTED MANUSCRIPT In our case, there are few dislocations, approximately 1.1–2.6 in a grain; that is, the dislocation pile-up effect represents only a small fraction of the material deformation
T
mechanism. The grains rotate or the grain boundaries slide instead of dislocations piling up.
IP
This rotation or sliding dominates the deformation mechanism. Consequently, specimens
SC R
B40–B150 were divided into two groups for discussion. One group comprises those specimens whose grain size is smaller than 10 nm; the other group comprises those whose grain size is larger than 10 nm. In the case of the small-grain-size (< 10 nm) group, grain
NU
rotation almost entirely dominates the deformation mechanism because the number of
MA
dislocations in a grain did not exceed two. Compared with the other specimens, specimens B75, B85, and B95 exhibit higher hardness, possibly because they simultaneously possess high compressive stress and a small grain size. The high compressive stress results in the
TE
D
grains becoming “clamped” during grain rotation. Under this condition, the increase of the frictional force is caused by the increase of the surface area of contact and by the difference
CE P
in the atomic sizes of Zr and Ti, which makes the contraction point uneven. The frictional force plays an important role in grain rotation. Therefore, rotation of grain with a high
AC
compressive stress and small grain size is difficult, further enhancing the hardness of the TiZrN films.
In the other case of grain sizes larger than 10 nm, although more dislocations exist in a grain, the number of dislocations in a grain is still very small. Therefore, we speculate that, in this case, few of the dislocations undergo deformation. Therefore, the hardness of the TiZrN films decreases as dislocations undergo deformation. Thus, the specimens with large grain sizes (> 10 nm) exhibit lower hardness compared to specimens B75, B85, and B95. In summary, the variation of residual stress was substantial as the substrate bias was varied from −35 V to −65 V. Therefore, the residual stress can be controlled to a lower value through adjustment of the substrate bias. At substrate biases of −40 V and −45 V, the residual
ACCEPTED MANUSCRIPT compressive stresses were 2.7–3.7 GPa. In addition, the films retained their hardness of 33.4–34.5 GPa.
T
Coloration is an important property in the decoration industry. Because TiZrN films
IP
exhibit a bright-golden color, they are widely used as decorative protective films. All the
SC R
optical data, including the brilliance L* and the a* and b* values, are listed in Table 2. The L* values for all samples are greater than 80, except in the case of sample B35. The B35 specimen exhibits lower brilliance, which may be a consequence of its larger roughness of
NU
approximately 4.4 nm. Incident light is easily absorbed and reflected at the surface of films
MA
with high roughness. Therefore, sample B35, which has a rough surface, appears dark. In total, the experimental results indicate that the window of processing parameters is
D
very wide for substrate biases in the range from −40 V to −120 V. The TiZrN films maintain
TE
excellent properties, including high hardness, brilliance, and low resistivity, when deposited at substrate biases within this range. Such a broad range of acceptable substrate biases is
AC
Conclusions
CE P
beneficial for industrial applications.
The window of processing parameters for TiZrN films is very wide in the substrate bias range from −40 V to −120 V. In this range, the TiZrN films exhibit high hardness (33.4 ~ 37.8 GPa), low resistivity (30 ~ 37 μΩ·cm), good roughness (0.5 ~ 0.6 nm), and a bright-golden color (brilliances greater than 80). At low substrate biases of −40 V and −45 V, the resulting TiZrN thin films retain both their high hardness (33.4 ~ 34.5 GPa) and their low residual compressive stress (2.7 ~ 3.7 GPa). The residual stress of TiZrN films gradually increases from −0.8 GPa to −5.5 GPa as the substrate bias is increased from −35 V to −65 V. At substrate biases more negative than −65 V, the residual stress ranges from −4.8 to −5.1 GPa. In the case of high residual stress and small grain size, grain rotation becomes difficult,
ACCEPTED MANUSCRIPT which further enhances the hardness of the TiZrN films. The hardness of the TiZrN films reached a maximum value of approximately 37.5 GPa. The resistivity of the TiZrN thin film
T
is related to its packing factor and the presence of lattice defects. The resistivity decreases
SC R
IP
with increasing packing factor.
Acknowledgment
This research was supported by the National Science Council of the Republic of China
NU
under contracts NSC 103-2221-E-007-101. The RBS analysis was carried out at the
MA
Accelerator Laboratory, National Tsing Hua University, Taiwan, R.O.C. The XPS analysis was carried out at the Instrument Center of the Material Science and Engineering Department,
D
National Tsing Hua University, Taiwan, R.O.C. SEM and XRD were performed in the
AC
CE P
TE
Instrument Center, National Chiao Tung University, Taiwan, R.O.C.
ACCEPTED MANUSCRIPT References
AC
CE P
TE
D
MA
NU
SC R
IP
T
[1] U. Wiiala, I. Penttinen, A. Korhonen, J. Aromaa, E. Ristolainen, Improved corrosion resistance of physical vapour deposition coated TiN and ZrN, Surf. Coat. Technol., 41 (1990) 191-204. [2] H. Holleck, Material selection for hard coatings, J. Vac. Sci. Technol., A, 4 (1986) 2661-2669. [3] W. D. Sproul, Very high-rate reactive sputtering of TiN, ZrN and HfN, Thin Solid Films 107 (1983) 141-147. [4] E. Török, A. Perry, L. Chollet, W. Sproul, Young's modulus of TiN, TiC, ZrN and HfN, Thin Solid Films 153 (1987) 37-43. [5] Y. W. Lin, J. H. Huang, G. P. Yu, Effect of nitrogen flow rate on properties of nanostructured TiZrN thin films produced by radio frequency magnetron sputtering, Thin Solid Films 518 (2010) 7308-7311. [6] Y. W. Lin, J. H. Huang, G. P. Yu, Microstructure and corrosion resistance of nanocrystalline TiZrN films on AISI 304 stainless steel substrate, J. Vac. Sci. Technol., A, 28 (2010) 774-778. [7] Y. W. Lin, J. H. Huang, G. P. Yu, C. N. Hsiao, F. Z. Chen., Influence of ion bombardment on structure and properties of TiZrN thin film, Appl. Surf. Sci. , 354 (2015) 155-160. [8] M. Ahlgren, H. Blomqvist, Influence of bias variation on residual stress and texture in TiAlN PVD coatings, Surf. Coat. Technol., 200 (2005) 157-160. [9] G. Taguchi, System of Experimental Design: Engineering Methods to Optimize Quality and Minimize Costs, 1, UNIPUB/Kraus International Publications, 1987. [10] G. Abadias, Y. Tse, P. Guerin, V. Pelosin, Interdependence between stress, preferred orientation, and surface morphology of nanocrystalline TiN thin films deposited by dual ion beam sputtering, J. Appl. Phys., 99 (2006) 113519-1-113519-13.. [11] W.-J. Chou, G.-P. Yu, J.-H. Huang, Bias effect of ion-plated zirconium nitride film on Si(100), Thin Solid Films 405 (2002) 162-169. [12] W.-J. Chou, G.-P. Yu, J.-H. Huang, Deposition of TiN thin films on Si(100) by HCD ion plating, Surf. Coat. Technol., 140 (2001) 206-214. [13] J Musil, J Vlček, Magnetron sputtering of hard nanocomposite coatings and their properties, Surf. Coat. Technol., 142-144 (2001) 557-566. [15] C.-H. Ma, J.-H. Huang, H. Chen, Nanohardness of nanocrystalline TiN thin films, Surf. Coat. Technol., 200 (2006) 3868-3875.
ACCEPTED MANUSCRIPT Table 1 Summary of the experimental results.
T
Table 2 Summary of the experimental results.
IP
Fig. 1 SEM cross-sectional images of TiZrN films deposited at a negative substrate
SC R
bias of 40 V.
Fig. 2 Deposition rate of TiZrN films as a function of the negative substrate bias.
NU
Fig. 3 Packing factor of TiZrN films as a function of the negative substrate bias.
MA
Fig. 4 SEM surface image of TiZrN films deposited at different substrate biases: samples B35, B40, B85, and B150.
TE
D
Fig. 5 AFM 3D surface images of samples B35, B40, B85, and B150.
CE P
Fig. 6 SEM image of a region of the B150 specimen that was delaminated by high-energy ion bombardment, which decreased the adhesion strength.
AC
Fig. 7 Variation of the hardness with the annealing temperature for TiZrN films with different compositions. Fig. 8 FWHM (111) of TiZrN films with respect to the negative substrate bias. Fig. 9 Grain size of TiZrN films with respect to the negative substrate bias. Fig. 10 Residual stress of TiZrN films with respect to the negative substrate bias. Fig. 11 Electrical resistivity of TiZrN films with respect to the negative substrate bias. Fig. 12 Resistivity of TiZrN films with respect to the packing factor. Fig. 13 Hardness of TiZrN films with respect to the negative substrate bias.
Fig.1
AC
CE P
TE
D
MA
NU
SC R
IP
T
ACCEPTED MANUSCRIPT
ACCEPTED MANUSCRIPT
T
8.5
IP
8.0
7.5
7.0
6.5 40
60
80
100
120
140
AC
Fig.2
CE P
TE
D
MA
Negative bias (V)
160
NU
20
SC R
Deposition rate (nm/min)
9.0
ACCEPTED MANUSCRIPT
0.98 0.96 0.94
T IP
0.90 0.88 0.86 0.84 0.82 0.80 40
60
80
100
120
TE
D
MA
Negative bias (V)
AC
CE P
Fig.3
140
NU
20
SC R
Packing factor
0.92
160
ACCEPTED MANUSCRIPT
AC
CE P
TE
D
MA
NU
SC R
IP
T
Figure 4
ACCEPTED MANUSCRIPT
AC
CE P
TE
D
MA
NU
SC R
IP
T
Figure 5
AC
CE P
TE
D
MA
NU
SC R
IP
T
ACCEPTED MANUSCRIPT
Fig. 6
TiN(200)
B150
IP
B120 Si(100)
B95
SC R
Intensuty (arb. units)
T
TiN(111) ZrN(200)
ZrN(111)
ACCEPTED MANUSCRIPT
B85 B75
NU
B65
25
30
35
40
45
50
TE
20
D
MA
B55
AC
CE P
2(degree)
Fig. 7
B45 B40
B35 55
60
ACCEPTED MANUSCRIPT
1.2 1.1 1.0
T
0.8
IP
0.7 0.6 0.5 0.4 0.3 0.2
Region 2 medium bias
Region 1 low bias
0.1
Region 3 high bias
0.0 40
60
80
100
120
AC
CE P
TE
D
MA
Negative bias (V)
140
NU
20
SC R
FWHM-111 (degree)
0.9
Fig.8
160
ACCEPTED MANUSCRIPT
28
Region 2 medium bias
Region 1 low bias
26
Region 3 high bias
24 22
T
18
IP
16 14 12 10 8 6 4 2 0 40
60
80
100
120
CE P
TE
D
MA
Negative bias (V)
AC
140
NU
20
SC R
Grain size (nm)
20
Fig.9
160
ACCEPTED MANUSCRIPT
6
T IP Region 1 low bias
3
Region 2 medium bias
Region 3 high bias
NU
2
SC R
4
1
0 20
40
60
MA
Compressive Stress (GPa)
5
80
100
AC
CE P
TE
D
Negetive Bias (V)
Fig. 10
120
140
160
ACCEPTED MANUSCRIPT
T
55
Region 2 medium bias
IP
Region 1 low bias
45
Region 3 high bias
SC R
Resistivity (u -cm )
50
40
30
20
40
60
80
100
NU
35
120
AC
CE P
TE
D
MA
Negative Bias (V)
Fig. 11
140
160
ACCEPTED MANUSCRIPT
55
IP
T
45
35
30 0.80
0.82
0.84
0.86
0.88
0.90
0.92
0.94
0.96
CE P
TE
D
MA
NU
Packing factor
SC R
40
AC
Resistivity (u -cm )
50
Fig. 12
ACCEPTED MANUSCRIPT
40 38 36
T
32
IP
30 28 26 24 22 20 40
60
80
100
120
AC
CE P
TE
D
MA
Negetive Bias (V)
140
NU
20
SC R
Hardness (GPa)
34
Fig. 13
160
ACCEPTED MANUSCRIPT Table 1 Summary of the experimental results.
531
1.1
1.1
B40
40
502
1.0
B45
45
495
1.0
B55
55
478
1.1
B65
65
476
1.1
B75
75
474
B85
85
B95
95
B150
T
IP
0.52
0.43
0.77
4.43
1.1
0.5 0.50 8
0.37
0.82
4.42
0.9
0.5 0.50 8
0.39
0.90
4.43
7
0.5 0.55 7
0.43
0.93
4.43
0.9
0.5 0.50 7
0.45
0.94
4.44
1.1
0.9
0.5 0.52 5
1.05
0.95
4.45
469
1.0
1.1
0.5 0.54 9
1.00
0.94
4.45
469
1.0
1.1
0.5 0.54 9
0.90
1.00
4.45
120
467
1.0
1.0
0.5 0.53 9
0.59
1.00
4.43
150
418
1.0
0.9
0.5 0.58 9
0.56
1.00
4.45
CE P
TE
D
1.1
AC
B120
0.5
NU
35
MA
B35
Zr/(Ti+Zr) FWHM (111) Lattice ratio ( 2θ) Texture parameter XPS (111) coefficient (Å) RBS
SC R
Thickne N/(Ti+Zr) Negative ss Sample ratio bias (nm) No. XPS (V) (±10nm RBS )
Deposited condition ：Ar flow rate 50 sccm N2 flow rate 19 sccm Ti target gun current 0.27 A Zr target gun current 0.27 A Deposition time：60 min
ACCEPTED MANUSCRIPT Table 2 Summary of the experimental results.
35 20
22.3 ± 1.5
313.3 ± 9.6
－0.8
52 ± 0.6
4.4
3. 38. 5 3
0.82
71 7. 26. .5 0 7
B40
40 22
33.4 ± 1.4
366.9 ± 16.9
－2.7
37 ± 0.6
0.6
0. 9.0 5
0.87
81 4. 29. .9 3 3
B45
45 21
34.5 ± 0.8
365.6 ± 14.7
－3.7
34 ± 0.5
0.5
0. 5.1 4
0.91
82 4. 30. .2 5 1
B55
55 19
34.5 ± 1.3
353.9 ± 18.4
－4.0
31 ± 0.4
0.5
0. 6.5 4
0.89
83 3. 29. .2 6 7
B65
65 19
34.6 ±
347.9 ±
1.4
12.9
B75
75
8
37.2 ± 1.8
B85
85
8
B95
95
9
D
MA
NU
B35
SC R
IP
T
Gra Young’s Roughness Coloratio Resid Nega in modules Resisti (nm) Packin n Sam Hardne ual tive siz g (GPa) stres vity ple ss bias e (μΩ-c Rms( Rma Factor a No. (GPa) s Ra L* b* (V) (n m) Rq) * x (±10%) (GPa) m)
30 ± 0.5
0.5
0. 4
7.2
0.95
83 3. 29. .5 9 5
TE
－5.5
33 ± 0.3
37.8 ± 1.3
366.4 ± 8.4
－4.8
33 ± 0.8
0.6
0. 6.8 5
0.92
83 3. 28. .1 6 9
37.6 ± 1.3
370.5 ± 16.5
－4.8
34 ± 1.0
0.5
0. 5.3 4
0.93
83 3. 28. .0 8 6
B12 34.7 ± 120 14 0 0.9
337.4 ± 12.6
－4.7
36 ± 0.9
0.6
0. 6.7 5
0.88
83 3. 30. .2 4 0
B15 33.2 ± 150 15 0 1.3
343.6 ± 8.5
－5.1
41 ± 0.7
0.8
0. 8.4 6
0.86
81 4. 29. .9 0 1
AC
CE P
351.3 ± 13.1
－5.1
0.5
0. 3.9 4
0.95
82 4. 29. .6 0 1
Deposited condition ：Ar flow rate 50 sccm N2 flow rate 19 sccm Ti target gun current 0.27 A Zr target gun current 0.27 A Deposition time：60 min
ACCEPTED MANUSCRIPT Highlights
CE P
TE
D
MA
NU
SC R
IP
T
The hardness of TiZrN reaches 37.5GPa. TiZrN maintains low residual stress. Resistivity is inversely proportional to the packing density.
AC
S0040-6090(16)30172-9 doi: 10.1016/j.tsf.2016.05.021 TSF 35204
To appear in:
Thin Solid Films
Received date: Revised date: Accepted date:
10 November 2015 16 May 2016 16 May 2016
Please cite this article as: Yu-Wei Lin, Hsi-An Chen, Ge-Ping Yu, Jia-Hong Huang, Effect of Bias on the Structure and Properties of TiZrN Thin Films Deposited by Unbalanced Magnetron Sputtering, Thin Solid Films (2016), doi: 10.1016/j.tsf.2016.05.021
This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Effect of Bias on the Structure and Properties of TiZrN Thin Films Deposited by
1
SC R
IP
Yu-Wei Lin1,3, Hsi-An Chen 1, Ge-Ping Yu1,2,*, Jia-Hong Huang1
T
Unbalanced Magnetron Sputtering
Department of Engineering and System Science, National Tsing Hua University, Taiwan,
R.O.C.
Institute of Nuclear Engineering and Science, National Tsing Hua University, Taiwan,
NU
2
3
MA
R.O.C.
Instrument Technology Research Center, National Applied Research Laboratories, Taiwan,
Presenter: Ge-Ping Yu
*
TE
†
D
R.O.C
AC
CE P
Corresponding author’s e-mail: [email protected]
ACCEPTED MANUSCRIPT ABSTRACT The objective is to investigate the substrate bias effect on the structure and properties
T
of the TiZrN thin films. The TiZrN thin films were deposited by direct current unbalanced
IP
magnetron sputtering system with dual guns (Ti and Zr) targets onto Si (100) substrates at
SC R
different substrate bias ranging from －35 V to －150 V. Experimental results indicated that all the specimens have strong (111) texture in X-ray Diffraction patterns. In this study, we
NU
discovered a transition bias of －35 V, above which a significant improvement of properties was found, including high hardness, excellent brilliance, low resistivity and fine surface
MA
morphology. Within the bias range of －40 to －120 V, 1the hardness of TiZrN films is around 35.5 GPa, the resistivity is about 33.5 μΩ-cm, and the brilliance is larger than 80. The
D
roughness is between 0.5 nm and 0.6 nm. The TiZrN films maintain excellent properties
TE
through a large range of applying bias, indicating that the process window is considerable wide. However, structure damage and thin film delamination were found when substrate bias
CE P
reached －150 V. The Rutherford backscatterng spectrometry result and scanning electron microscope image further support the structure damage at －150 V. For protective coatings,
AC
low residual stress is required to avoid delamination. By adjusting substrate bias, residual stress can be controlled to lower value. In this study, the residual stress of TiZrN films gradually decreases with decreasing the substrate bias ranging from －65 V to －35 V. The TiZrN thin films with high hardness, lower residual stress could be obtained simultaneously at low substrate bias of －40 V and －45 V, the hardness and residual stress are 33.4～34.5 GPa and －2.7 ～－3.7 GPa, respectively. Keywords: TiZrN, hardness, residual stress
ACCEPTED MANUSCRIPT Introduction Owing to their excellent mechanical properties, low resistance, golden color, and
T
superior corrosion resistance, titanium nitride (TiN) and zirconium nitride (ZrN) have been
IP
widely used in various industries over the past 20 years[1-4]. Their applications include
SC R
decorative coatings, diffusion barriers in microelectromechanical systems, and hard coatings on tools. However, the demand for harder and tougher materials driven by technological progress necessitates the development of ternary transition-metal nitride films such as TiZrN,
NU
TiSiN, and TiAlN. Compared with TiN and ZrN, ternary TiZrN films not only exhibit a low
MA
resistivity and golden color but also feature superior mechanical properties and better corrosion resistance[5, 6].
D
Previous study [7] indicated that the addition of Zr into TiN or Ti into ZrN could
TE
enhance the fracture toughness of hard coatings. This approach of incorporating a third element is effective for improving various properties of thin films. In general, the high
CE P
hardness of hard coatings usually accompanies high residual stress in films deposited via the physical vapor deposition (PVD) process. However, the presence of such high residual stress
AC
in thin films can result in reduced adhesion and film spallation. The residual stress can be reduced while retaining the excellent mechanical properties of films by controlling the substrate bias during deposition[8]. The substrate bias is an important parameter in the PVD process. It governs the total energy and momentum of adatoms while influencing the structure and properties of the resulting thin films. In addition, the substrate bias is easily adjusted during thin film deposition. Therefore, we chose to change the substrate bias as a single-variable experimental parameter. Our objective was to investigate the effect of substrate bias on the structure and properties of deposited TiZrN thin films. The influence of substrate bias on the microstructure, surface morphology, hardness, residual stress, resistivity, and packing factor
ACCEPTED MANUSCRIPT was investigated. Experiments
T
In this research, ternary TiZrN thin films were deposited using a reactive direct current
IP
unbalanced magnetron (UBM) sputtering sputtering system. p-type Si(100) wafers were
SC R
chosen as the substrate material; the dimensions of the Si substrates were 3.5 × 3.5 × 0.525 cm3. The sputtering sources were a Zr target (99.9%) and a Ti target (99.995%), each with a diameter of 2 inches.
NU
First, before the coating process, the specimens were ultrasonically cleaned sequentially in acetone and methanol for 5 min each. The chamber was evacuated to a base pressure of 6.7
MA
× 10−4 Pa using a pumping system at 400°C. The substrate was then pre-sputtered with Ar+
D
ions at a bias of −1000 V for 5 min with the Ar (99.9995% purity) flow rate fixed at 50 sccm
TE
to remove the surface oxide layer.
After pre-sputtering, the working gas (Ar) and reactive gas (N2) were introduced into the
CE P
chamber. The Ar gas flow rate was fixed at 30 sccm with a maximum gas flow rate of 50 sccm; the N2 gas flow rate was fixed at 1.9 sccm with a maximum gas flow rate of 5 sccm.
AC
The total gas pressure was maintained at 1.5 × 10−1 Pa. The coating temperature was 400°C and was monitored using a thermocouple positioned between the substrate holder and the halogen lamps. During deposition, the dual gun current was kept constant at 0.27 A. The sample number was assigned according to the substrate bias (e.g., sample B35 = 35 V substrate negative bias). The substrate bias was varied from −35 V to −150 V. The target-to-substrate distance was adjusted to 10 cm. The TiZrN thin films were deposited by co-sputtering of Ti and Zr targets. The specimens were rotated during sputtering. Both normal lines of the Ti and Zr target surfaces intersected on the revolved substrate surface to ensure chemical uniformity of the coating. The deposition time of all specimens was 60 min. The optimum parameters were determined by the Taguchi method[9].
ACCEPTED MANUSCRIPT All the deposited specimens were analyzed. The surface compositions of TiZrN films were determined with a PHI 1600 X-ray photoelectron spectroscope at a vacuum of 10-8 Pa.
T
The X-ray source was Mg Kα (1253.6 eV). In order to remove the surface oxide, prior to
IP
analysis, the specimen surface was pre-sputtered by 3 keVAr+ for 3 minutes. The
SC R
deconvolution and fitting of X-ray photoelectron spectroscopy (XPS) spectrum was executed by a computer program XPSPEAK 4.1. The background of the spectra was subtracted by Shirley method. The high resolution spectra were then deconvoluted and processed by
NU
least-square fitting procedure using Gaussian function. The atomic percentages of Ti, Zr, N
MA
and O were obtained from the integrated areas of deconvoluted spectra divided by the corresponding sensitivity factors. The relative content of each element was calculated by the 100%, where i represents Ti, Zr, N, or O, Ci is the content of
D
following equation: Ci=
TE
element i in the thin films, and I is the integrated intensity of the element i divided by the corresponding sensitivity factor. The X-ray source was Mg Kα (1253.6 eV). The N/(Ti+Zr)
CE P
ratio and packing factor were determined by Rutherford backscattering spectrometry (RBS). The Van de Graaff accelerator at NTHU was applied. The incident ion source was α (4He+)
AC
particles and the operating energy was 2 MeV. The chamber was evacuated below 8 × 10-4 Pa. A surface barrier detector was set up at the backscattering angle of 160° to the surface normal of specimens. Prior to analysis, the RBS was calibrated using Si and W standard samples. The resolution of the instrument was about 30 ~100 Å. The crystal structure and preferred orientation of TiZrN films were characterized by X-ray diffraction (XRD) using a D8 diffractometer. The Cu Kα line at 1.5405 Å was used as the source for diffraction. For the θ/2θ scans, a scan rate of 4deg/min was used, and the 2θ scan range was from 20 to 60
I ( hkl ) degrees. The texture coefficient of (hkl) orientation was defined as I ( hkl ) , where I is the integrated intensity of the corresponding (hkl) peak. The thickness and morphology of the
ACCEPTED MANUSCRIPT TiZrN thin films were observed by a JOEL JSM-6700F scanning electron microscopy (SEM) at National Chung Hsing University. The cross-sectional microstructure of TiZrN/Si was also
T
observed in SEM image. Before loading, the TiZrN/Si samples were cleaved into a piece of
IP
0.3×0.3 cm2 and attached on a sample holder with carbon tape and then the samples were
SC R
coated an Au layer to improve electric conductivity. The surface roughness and morphology of TiZrN thin films were observed by an atomic force microscope (AFM) with a silicon nitride tip. During the analysis, the contact mode which means that a constant contact force
NU
applied upon the specimen surface was chosen. A scanned frequency of 4 Hz line by line and
MA
2 × 2 μm2 scan size were set. The residual stress in the TiZrN thin films was measured using the laser curvature method. The electrical resistivity of the films was measured by the
D
four-point probe method. The coloration of the TiZrN coatings was characterized at room
TE
temperature using a Hunter Lab Miniscan XE Plus spectrophotometer, model 4000VSAV. The hardness of the deposited films was determined by nanoindentation tests
CE P
performed using a triangular Berkovich indenter. Sample hardness and elastic modulus were inferred from the results of the indentation tests. The tip was moved laterally across a surface
AC
while the normal force was varied or held constant. Displacement into the surface and lateral (tangential) force were measured during the process. The film hardness is theoretically a constant value but will vary with changes in the applied load. In general, the indentation depth should be less than one tenth of the film thickness to avoid the substrate effect. To evaluate the creep effect, the load was first increased over a period of 20 s to the desired load, maintained for 5 s, and then unloaded over a period of 20 s. The unloading stiffness was interpreted by the equation
S
dP 2 Er A dh
ACCEPTED MANUSCRIPT where the unloading stiffness S = dP/dh is measured experimentally, A is the projected area of the elastic contact, and Er is the reduced modulus.
IP
SC R
1 (1 2 ) (1 i2 ) Er E Ei
T
The reduced modulus is defined as
where E and Ei are the Young’s moduli for the specimen and indenter, respectively, and ν and νi are the Poisson’s ratios for the specimen and indenter, respectively. The reduced modulus
NU
can be determined if the contact area can be independently measured. Parameter A is a
MA
function of the contact depth (hc), and hc can be calculated by the following equation
hc hmax hs
D
where hmax is the total displacement determined from experimental data and hs is the depth of
TE
the surface at the surrounding of the contact:
( 2)
CE P
hs
(h h f )
2( 2) Pmax S
AC
where hf is the final depth of after load releasing and Pmax is the maximum load. Thus,
hc hmax 0.72
Pmax S
Finally, the hardness can also be determined as
H
Pmax A
where H is the film’s hardness, Pmax is the maximum load, and A is the projected area of contact. Before measurement of the coating samples, a quartz sample should be used to calibrate the Cx value of A:
A(hc ) C0 hc2 C1hc1 C2 hc1 / 2 C3 hc1 / 4 C4 hc1 / 6 C5 hc1 / 8
ACCEPTED MANUSCRIPT where C0 is a constant value of 24.5 for the Berkovich indenter, and C1–C5 values can be fitted using five different loads and contact depths. Notably, the hardness calibration value
T
and the reduced modulus of the quartz should be approximately 10 (±3%) GPa and 69.6
SC R
IP
(±3%) GPa, respectively.
Results and Discussion
According to XPS and RBS analyses, the N/(Ti+Zr) ratio for all samples is fixed at
NU
0.9–1.1, and the Zr/(Ti+Zr) ratios are approximately 0.55 (Table 1). The chemical
MA
composition of TiZrN thin films is not sensitive to substrate bias. An increase of the substrate bias not only increases the total energy but also enhances the momentum of the adatoms. The
D
momentum of an adatom is directly proportional to its mobility. Adatoms can move to stable
TE
sites as they obtain sufficient momentum, resulting in deposition of a denser thin film. By contrast, an excessively high negative substrate bias results in negative influences such as the
CE P
production of additional lattice defects, radiation damage, and thin film delamination. Fig. 1 presents SEM cross-sectional images showing the microstructure and thickness of
AC
the B40 specimens. The thickness of the TiZrN films ranges from 418 to 531 nm. The deposition rate was calculated as the thickness divided by the deposition time. The results demonstrate that the deposition rate decreases with increasing substrate bias because of the re-sputtering effect. The variation of deposition rate under different substrate biases is plotted in Fig. 2. Fig. 3 reveals that the packing factor of the films gradually increases from 0.82 to 0.96 when the substrate bias is between −35 V and −65 V; it then slightly decreases when the substrate bias exceeds −65 V. The aforementioned SEM cross-sectional results also correspond well with the packing factor. Fig. 3 also shows that the B35 and B150 specimens have lower packing factors.
ACCEPTED MANUSCRIPT Figs. 4 and 5 show the SEM surface images and AFM 3D images. The roughness of all samples is also given in Table 2. When the substrate bias exceeds −35 V, the adatoms obtain
T
sufficient momentum to move to stable sites; consequently, the TiZrN films become very
IP
smooth. Their roughness ranges from 0.5 nm to 0.8 nm. Although sample B150 also exhibits
SC R
a flat surface, a region of the B150 specimen delaminated because of high-energy ion bombardment, decreasing the adhesion strength (Fig. 6).
Fig. 7 shows XRD patterns of the TiZrN films deposited under different substrate biases.
NU
In XRD patterns, the reflection of the TiZrN (111) peak is observed between those of the TiN
MA
(111) peak and the ZrN (111) peak; similarly, the reflection of the TiZrN (200) peak is observed between those of the TiN (200) peak and the ZrN (200) peak. Moreover, according
D
to Hume–Rothery rules, TiZrN should be a substitutional solid solution. The difference
TE
between the diameters of the atoms (Zr and Ti) is less than 15%, the crystal structures of TiN and ZrN are both face-centered cubic, the electronegativities of Ti and Zn do not substantially
CE P
differ, and they both belong to the same group (IVB) in the periodic table. Thus, TiN and ZrN satisfy the generalized definition for a substitutional solid solution. The texture coefficient
AC
can be obtained by ITiZrN(111)/(ITiZrN(111) + ITiZrN(200)), where I(111) and I(200) represent the integrated intensities of the corresponding (111) and (200) peaks, respectively. Actually, all the specimens have strong (111) texture that is greater than or equal to 0.77. The B35 specimen has a lower (111) texture coefficient at −35 V, possibly because its looser structure results in random crystalline orientation. Therefore, the (111) texture coefficients increase with increasing substrate bias. Furthermore, Abadias proposed that texture variation can be attributed to a re-nucleation mechanism[10]. A high substrate bias results in the formation of more defects in the resulting thin films, and regions with a high defect density are prone to secondary nucleation and grain-boundary migration. Therefore, more grains grew along the (111) direction in the films grown in this study.
ACCEPTED MANUSCRIPT The full-width at half-maximum (FWHM) of TiZrN (111) as a function of substrate bias is shown in Fig. 8, where a smaller FWHM indicates good crystallinity of the TiZrN film. Fig.
T
9 illustrates variation of grain size as a function of substrate bias. The vertical dimension of
IP
the grains in TiZrN can be calculated using the Scherrer equation. The grain size of TiZrN
SC R
films for all samples is less than 22 nm. The FWHM(111) and grain size results are summarized in Table 2. The B35–B150 specimens were divided into three regions for purposes of discussion:
NU
Region 1: Low-negative-bias region with substrate bias ranging from −35 V to −65 V Region 2: Medium-negative-bias region with substrate bias ranging from −75 V to −95 V
MA
Region 3: High-negative-bias region with substrate bias ranging from −120 V to −150 V
Originally, the large grain size was approximately 20 nm in region 1. As the substrate
D
bias exceeded −65 V (entering region 2), the grain size decreased to approximately 8.3 nm.
TE
This decrease in grain size is attributed to the presence of additional defects induced by ion
CE P
bombardment. Areas with a high defect density are prone to nucleation; consequently, nucleation increases simultaneously with increasing number of defects, which decreases the grain size. The FWHM of (111) suddenly increased by 0.6° because the number of defects
AC
increased when the substrate bias exceeded −65 V. In region 3, the FWHM of (111) slightly decreased, possibly because the grains lacked sufficient energy for rearrangement of the (111) orientations at high substrate bias. Thus, the crystallinity of the films was enhanced. We observed similar behavior in our previous study[11]. According to the Scherrer equation, the grain size in the present study increased to approximately 14.5 nm from 8.3 nm. Compared to TiN and ZrN films, TiZrN films possess higher residual stress, which is related to the difference between the atomic radii: the atomic radius of Zr is larger than that of Ti. The formation of a substitutional solid solution by the addition of Zr into the TiN lattice or by the addition of Ti into the ZrN lattice induces serious lattice distortion, resulting in high residual stress. Fig. 10 shows the variation of the residual stress as a function of
ACCEPTED MANUSCRIPT substrate bias. It shows that the residual stress gradually increases from −0.8 to −5.5 GPa in region 1. The impact of high-energy Ar+ ions on a film surface during the deposition process
T
increases the residual stress in the film. The residual stress slightly decreased with increasing
IP
substrate bias in region 2, where the increase in the mobility of adatoms to stable sites
SC R
released the residual stress. When the substrate bias exceeded −95 V (region 3), the residual stress slightly increased because additional lattice defects formed under the high negative bias.
NU
Fig. 11 presents the variation of the electrical resistivity as a function of the substrate
MA
bias. The results show that all the samples exhibited low resistivity (from 30 to 52 μΩ·cm). Initially, the electrical resistivity suddenly decreased when the substrate bias exceeded −35 V in region 1. It thereafter remained constant (from 30 to 37 μΩ·cm) as the substrate bias was
TE
D
increased from −40 to −120 V. The electrical resistivity then slightly increased when the substrate bias exceeded −150 V.
CE P
The resistivity of the films is related to their number of lattice defects. An increase in the number of lattice defects can impede the movement of free electrons and increase their mean
AC
free path. Consequently, more lattice defects exist in the film, which may increase the film’s resistivity. The resistivity rapidly decreased when the substrate bias exceeded −35 V, which we attribute to the increase in densification. Increasing densification of the film results in a decrease in the number of lattice defects, further decreasing the resistivity of the thin films. However, in the case where the substrate bias was more negative than −95 V (region 3), the corresponding increase in resistivity is attributed to the fact that high-energy ion bombardment induced structural damage, resulting in a slight increase in the resistivity of the TiZrN films because of the introduction of additional defects. Therefore, we plotted the resistivity as a function of the packing factor in Fig. 12. The resulting plot indicates that the lowest packing corresponds to the highest resistivity.
ACCEPTED MANUSCRIPT Fig. 13 shows that the hardness of all specimens varied with the substrate bias. The hardness values of all the TiZrN films exceeded 33 GPa, ranging from 33.2 to 37.8 GPa,
T
except in the case of sample B35. Sample B35 exhibited the lowest hardness because of its
IP
looser structure, consistent with the SEM cross-section analysis and RBS results. The reason
SC R
the hardness of TiZrN films exceeds 33 GPa is likely associated with solid-solution strengthening. The addition of Zr into the TiN lattice or the addition of Ti into the ZrN lattice induces lattice distortion. The lattice distortion impedes the movement of dislocations and
NU
grain boundaries, thereby enhancing the hardness of TiZrN films.
MA
Moreover, texture, residual stress, grain size, and packing factor also influence the hardness of the TiZrN films. In NaCl-type TiZrN, the {110} planes and <110> direction
D
represent an active slip system. If the external force is parallel to the (111) plane normal, the
TE
Schmid factor is equal to zero on all slip systems[12]. Therefore, <111> is the hardest orientation, and the hardness of the film is enhanced with increasing (111) texture coefficient.
CE P
Similar conditions have been observed in the TiN system[13]. However, in our case, the hardness of the TiZrN films is not sensitive to changes in the (111) texture coefficient. The
AC
hardness of the TiZrN films might be affected by other factors such as residual stress, grain size, and packing factor. In this research, the grain size ranges from 8 nm to 21 nm for all samples. According to dislocation theory[14], the number of dislocations in a grain can be calculated; the authors of a previous study proposed a similar calculation for TiN thin films[15]. We can use the following equation to calculate the number of dislocations n in a grain:
where k is a factor that approaches unity, τ is the critical resolved shear stress, D is the grain size, G is the shear modulus, and b is the Burgers vector.
ACCEPTED MANUSCRIPT In our case, there are few dislocations, approximately 1.1–2.6 in a grain; that is, the dislocation pile-up effect represents only a small fraction of the material deformation
T
mechanism. The grains rotate or the grain boundaries slide instead of dislocations piling up.
IP
This rotation or sliding dominates the deformation mechanism. Consequently, specimens
SC R
B40–B150 were divided into two groups for discussion. One group comprises those specimens whose grain size is smaller than 10 nm; the other group comprises those whose grain size is larger than 10 nm. In the case of the small-grain-size (< 10 nm) group, grain
NU
rotation almost entirely dominates the deformation mechanism because the number of
MA
dislocations in a grain did not exceed two. Compared with the other specimens, specimens B75, B85, and B95 exhibit higher hardness, possibly because they simultaneously possess high compressive stress and a small grain size. The high compressive stress results in the
TE
D
grains becoming “clamped” during grain rotation. Under this condition, the increase of the frictional force is caused by the increase of the surface area of contact and by the difference
CE P
in the atomic sizes of Zr and Ti, which makes the contraction point uneven. The frictional force plays an important role in grain rotation. Therefore, rotation of grain with a high
AC
compressive stress and small grain size is difficult, further enhancing the hardness of the TiZrN films.
In the other case of grain sizes larger than 10 nm, although more dislocations exist in a grain, the number of dislocations in a grain is still very small. Therefore, we speculate that, in this case, few of the dislocations undergo deformation. Therefore, the hardness of the TiZrN films decreases as dislocations undergo deformation. Thus, the specimens with large grain sizes (> 10 nm) exhibit lower hardness compared to specimens B75, B85, and B95. In summary, the variation of residual stress was substantial as the substrate bias was varied from −35 V to −65 V. Therefore, the residual stress can be controlled to a lower value through adjustment of the substrate bias. At substrate biases of −40 V and −45 V, the residual
ACCEPTED MANUSCRIPT compressive stresses were 2.7–3.7 GPa. In addition, the films retained their hardness of 33.4–34.5 GPa.
T
Coloration is an important property in the decoration industry. Because TiZrN films
IP
exhibit a bright-golden color, they are widely used as decorative protective films. All the
SC R
optical data, including the brilliance L* and the a* and b* values, are listed in Table 2. The L* values for all samples are greater than 80, except in the case of sample B35. The B35 specimen exhibits lower brilliance, which may be a consequence of its larger roughness of
NU
approximately 4.4 nm. Incident light is easily absorbed and reflected at the surface of films
MA
with high roughness. Therefore, sample B35, which has a rough surface, appears dark. In total, the experimental results indicate that the window of processing parameters is
D
very wide for substrate biases in the range from −40 V to −120 V. The TiZrN films maintain
TE
excellent properties, including high hardness, brilliance, and low resistivity, when deposited at substrate biases within this range. Such a broad range of acceptable substrate biases is
AC
Conclusions
CE P
beneficial for industrial applications.
The window of processing parameters for TiZrN films is very wide in the substrate bias range from −40 V to −120 V. In this range, the TiZrN films exhibit high hardness (33.4 ~ 37.8 GPa), low resistivity (30 ~ 37 μΩ·cm), good roughness (0.5 ~ 0.6 nm), and a bright-golden color (brilliances greater than 80). At low substrate biases of −40 V and −45 V, the resulting TiZrN thin films retain both their high hardness (33.4 ~ 34.5 GPa) and their low residual compressive stress (2.7 ~ 3.7 GPa). The residual stress of TiZrN films gradually increases from −0.8 GPa to −5.5 GPa as the substrate bias is increased from −35 V to −65 V. At substrate biases more negative than −65 V, the residual stress ranges from −4.8 to −5.1 GPa. In the case of high residual stress and small grain size, grain rotation becomes difficult,
ACCEPTED MANUSCRIPT which further enhances the hardness of the TiZrN films. The hardness of the TiZrN films reached a maximum value of approximately 37.5 GPa. The resistivity of the TiZrN thin film
T
is related to its packing factor and the presence of lattice defects. The resistivity decreases
SC R
IP
with increasing packing factor.
Acknowledgment
This research was supported by the National Science Council of the Republic of China
NU
under contracts NSC 103-2221-E-007-101. The RBS analysis was carried out at the
MA
Accelerator Laboratory, National Tsing Hua University, Taiwan, R.O.C. The XPS analysis was carried out at the Instrument Center of the Material Science and Engineering Department,
D
National Tsing Hua University, Taiwan, R.O.C. SEM and XRD were performed in the
AC
CE P
TE
Instrument Center, National Chiao Tung University, Taiwan, R.O.C.
ACCEPTED MANUSCRIPT References
AC
CE P
TE
D
MA
NU
SC R
IP
T
[1] U. Wiiala, I. Penttinen, A. Korhonen, J. Aromaa, E. Ristolainen, Improved corrosion resistance of physical vapour deposition coated TiN and ZrN, Surf. Coat. Technol., 41 (1990) 191-204. [2] H. Holleck, Material selection for hard coatings, J. Vac. Sci. Technol., A, 4 (1986) 2661-2669. [3] W. D. Sproul, Very high-rate reactive sputtering of TiN, ZrN and HfN, Thin Solid Films 107 (1983) 141-147. [4] E. Török, A. Perry, L. Chollet, W. Sproul, Young's modulus of TiN, TiC, ZrN and HfN, Thin Solid Films 153 (1987) 37-43. [5] Y. W. Lin, J. H. Huang, G. P. Yu, Effect of nitrogen flow rate on properties of nanostructured TiZrN thin films produced by radio frequency magnetron sputtering, Thin Solid Films 518 (2010) 7308-7311. [6] Y. W. Lin, J. H. Huang, G. P. Yu, Microstructure and corrosion resistance of nanocrystalline TiZrN films on AISI 304 stainless steel substrate, J. Vac. Sci. Technol., A, 28 (2010) 774-778. [7] Y. W. Lin, J. H. Huang, G. P. Yu, C. N. Hsiao, F. Z. Chen., Influence of ion bombardment on structure and properties of TiZrN thin film, Appl. Surf. Sci. , 354 (2015) 155-160. [8] M. Ahlgren, H. Blomqvist, Influence of bias variation on residual stress and texture in TiAlN PVD coatings, Surf. Coat. Technol., 200 (2005) 157-160. [9] G. Taguchi, System of Experimental Design: Engineering Methods to Optimize Quality and Minimize Costs, 1, UNIPUB/Kraus International Publications, 1987. [10] G. Abadias, Y. Tse, P. Guerin, V. Pelosin, Interdependence between stress, preferred orientation, and surface morphology of nanocrystalline TiN thin films deposited by dual ion beam sputtering, J. Appl. Phys., 99 (2006) 113519-1-113519-13.. [11] W.-J. Chou, G.-P. Yu, J.-H. Huang, Bias effect of ion-plated zirconium nitride film on Si(100), Thin Solid Films 405 (2002) 162-169. [12] W.-J. Chou, G.-P. Yu, J.-H. Huang, Deposition of TiN thin films on Si(100) by HCD ion plating, Surf. Coat. Technol., 140 (2001) 206-214. [13] J Musil, J Vlček, Magnetron sputtering of hard nanocomposite coatings and their properties, Surf. Coat. Technol., 142-144 (2001) 557-566. [15] C.-H. Ma, J.-H. Huang, H. Chen, Nanohardness of nanocrystalline TiN thin films, Surf. Coat. Technol., 200 (2006) 3868-3875.
ACCEPTED MANUSCRIPT Table 1 Summary of the experimental results.
T
Table 2 Summary of the experimental results.
IP
Fig. 1 SEM cross-sectional images of TiZrN films deposited at a negative substrate
SC R
bias of 40 V.
Fig. 2 Deposition rate of TiZrN films as a function of the negative substrate bias.
NU
Fig. 3 Packing factor of TiZrN films as a function of the negative substrate bias.
MA
Fig. 4 SEM surface image of TiZrN films deposited at different substrate biases: samples B35, B40, B85, and B150.
TE
D
Fig. 5 AFM 3D surface images of samples B35, B40, B85, and B150.
CE P
Fig. 6 SEM image of a region of the B150 specimen that was delaminated by high-energy ion bombardment, which decreased the adhesion strength.
AC
Fig. 7 Variation of the hardness with the annealing temperature for TiZrN films with different compositions. Fig. 8 FWHM (111) of TiZrN films with respect to the negative substrate bias. Fig. 9 Grain size of TiZrN films with respect to the negative substrate bias. Fig. 10 Residual stress of TiZrN films with respect to the negative substrate bias. Fig. 11 Electrical resistivity of TiZrN films with respect to the negative substrate bias. Fig. 12 Resistivity of TiZrN films with respect to the packing factor. Fig. 13 Hardness of TiZrN films with respect to the negative substrate bias.
Fig.1
AC
CE P
TE
D
MA
NU
SC R
IP
T
ACCEPTED MANUSCRIPT
ACCEPTED MANUSCRIPT
T
8.5
IP
8.0
7.5
7.0
6.5 40
60
80
100
120
140
AC
Fig.2
CE P
TE
D
MA
Negative bias (V)
160
NU
20
SC R
Deposition rate (nm/min)
9.0
ACCEPTED MANUSCRIPT
0.98 0.96 0.94
T IP
0.90 0.88 0.86 0.84 0.82 0.80 40
60
80
100
120
TE
D
MA
Negative bias (V)
AC
CE P
Fig.3
140
NU
20
SC R
Packing factor
0.92
160
ACCEPTED MANUSCRIPT
AC
CE P
TE
D
MA
NU
SC R
IP
T
Figure 4
ACCEPTED MANUSCRIPT
AC
CE P
TE
D
MA
NU
SC R
IP
T
Figure 5
AC
CE P
TE
D
MA
NU
SC R
IP
T
ACCEPTED MANUSCRIPT
Fig. 6
TiN(200)
B150
IP
B120 Si(100)
B95
SC R
Intensuty (arb. units)
T
TiN(111) ZrN(200)
ZrN(111)
ACCEPTED MANUSCRIPT
B85 B75
NU
B65
25
30
35
40
45
50
TE
20
D
MA
B55
AC
CE P
2(degree)
Fig. 7
B45 B40
B35 55
60
ACCEPTED MANUSCRIPT
1.2 1.1 1.0
T
0.8
IP
0.7 0.6 0.5 0.4 0.3 0.2
Region 2 medium bias
Region 1 low bias
0.1
Region 3 high bias
0.0 40
60
80
100
120
AC
CE P
TE
D
MA
Negative bias (V)
140
NU
20
SC R
FWHM-111 (degree)
0.9
Fig.8
160
ACCEPTED MANUSCRIPT
28
Region 2 medium bias
Region 1 low bias
26
Region 3 high bias
24 22
T
18
IP
16 14 12 10 8 6 4 2 0 40
60
80
100
120
CE P
TE
D
MA
Negative bias (V)
AC
140
NU
20
SC R
Grain size (nm)
20
Fig.9
160
ACCEPTED MANUSCRIPT
6
T IP Region 1 low bias
3
Region 2 medium bias
Region 3 high bias
NU
2
SC R
4
1
0 20
40
60
MA
Compressive Stress (GPa)
5
80
100
AC
CE P
TE
D
Negetive Bias (V)
Fig. 10
120
140
160
ACCEPTED MANUSCRIPT
T
55
Region 2 medium bias
IP
Region 1 low bias
45
Region 3 high bias
SC R
Resistivity (u -cm )
50
40
30
20
40
60
80
100
NU
35
120
AC
CE P
TE
D
MA
Negative Bias (V)
Fig. 11
140
160
ACCEPTED MANUSCRIPT
55
IP
T
45
35
30 0.80
0.82
0.84
0.86
0.88
0.90
0.92
0.94
0.96
CE P
TE
D
MA
NU
Packing factor
SC R
40
AC
Resistivity (u -cm )
50
Fig. 12
ACCEPTED MANUSCRIPT
40 38 36
T
32
IP
30 28 26 24 22 20 40
60
80
100
120
AC
CE P
TE
D
MA
Negetive Bias (V)
140
NU
20
SC R
Hardness (GPa)
34
Fig. 13
160
ACCEPTED MANUSCRIPT Table 1 Summary of the experimental results.
531
1.1
1.1
B40
40
502
1.0
B45
45
495
1.0
B55
55
478
1.1
B65
65
476
1.1
B75
75
474
B85
85
B95
95
B150
T
IP
0.52
0.43
0.77
4.43
1.1
0.5 0.50 8
0.37
0.82
4.42
0.9
0.5 0.50 8
0.39
0.90
4.43
7
0.5 0.55 7
0.43
0.93
4.43
0.9
0.5 0.50 7
0.45
0.94
4.44
1.1
0.9
0.5 0.52 5
1.05
0.95
4.45
469
1.0
1.1
0.5 0.54 9
1.00
0.94
4.45
469
1.0
1.1
0.5 0.54 9
0.90
1.00
4.45
120
467
1.0
1.0
0.5 0.53 9
0.59
1.00
4.43
150
418
1.0
0.9
0.5 0.58 9
0.56
1.00
4.45
CE P
TE
D
1.1
AC
B120
0.5
NU
35
MA
B35
Zr/(Ti+Zr) FWHM (111) Lattice ratio ( 2θ) Texture parameter XPS (111) coefficient (Å) RBS
SC R
Thickne N/(Ti+Zr) Negative ss Sample ratio bias (nm) No. XPS (V) (±10nm RBS )
Deposited condition ：Ar flow rate 50 sccm N2 flow rate 19 sccm Ti target gun current 0.27 A Zr target gun current 0.27 A Deposition time：60 min
ACCEPTED MANUSCRIPT Table 2 Summary of the experimental results.
35 20
22.3 ± 1.5
313.3 ± 9.6
－0.8
52 ± 0.6
4.4
3. 38. 5 3
0.82
71 7. 26. .5 0 7
B40
40 22
33.4 ± 1.4
366.9 ± 16.9
－2.7
37 ± 0.6
0.6
0. 9.0 5
0.87
81 4. 29. .9 3 3
B45
45 21
34.5 ± 0.8
365.6 ± 14.7
－3.7
34 ± 0.5
0.5
0. 5.1 4
0.91
82 4. 30. .2 5 1
B55
55 19
34.5 ± 1.3
353.9 ± 18.4
－4.0
31 ± 0.4
0.5
0. 6.5 4
0.89
83 3. 29. .2 6 7
B65
65 19
34.6 ±
347.9 ±
1.4
12.9
B75
75
8
37.2 ± 1.8
B85
85
8
B95
95
9
D
MA
NU
B35
SC R
IP
T
Gra Young’s Roughness Coloratio Resid Nega in modules Resisti (nm) Packin n Sam Hardne ual tive siz g (GPa) stres vity ple ss bias e (μΩ-c Rms( Rma Factor a No. (GPa) s Ra L* b* (V) (n m) Rq) * x (±10%) (GPa) m)
30 ± 0.5
0.5
0. 4
7.2
0.95
83 3. 29. .5 9 5
TE
－5.5
33 ± 0.3
37.8 ± 1.3
366.4 ± 8.4
－4.8
33 ± 0.8
0.6
0. 6.8 5
0.92
83 3. 28. .1 6 9
37.6 ± 1.3
370.5 ± 16.5
－4.8
34 ± 1.0
0.5
0. 5.3 4
0.93
83 3. 28. .0 8 6
B12 34.7 ± 120 14 0 0.9
337.4 ± 12.6
－4.7
36 ± 0.9
0.6
0. 6.7 5
0.88
83 3. 30. .2 4 0
B15 33.2 ± 150 15 0 1.3
343.6 ± 8.5
－5.1
41 ± 0.7
0.8
0. 8.4 6
0.86
81 4. 29. .9 0 1
AC
CE P
351.3 ± 13.1
－5.1
0.5
0. 3.9 4
0.95
82 4. 29. .6 0 1
Deposited condition ：Ar flow rate 50 sccm N2 flow rate 19 sccm Ti target gun current 0.27 A Zr target gun current 0.27 A Deposition time：60 min
ACCEPTED MANUSCRIPT Highlights
CE P
TE
D
MA
NU
SC R
IP
T
The hardness of TiZrN reaches 37.5GPa. TiZrN maintains low residual stress. Resistivity is inversely proportional to the packing density.
AC