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Ti nitride phases in thin films deposited by DC magnetron sputtering
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surface science ELSEVIER
Applied SurfaceScience91 (1995)295-302
Ti nitride phases in thin films deposited by DC magnetron sputtering Rodica Manaila a, ,, Domokos Biro b, Peter B. Barna c, MiNos Adamik c, Florin Zavaliche a, Stefan Craciun d, Andrei Devenyi a a Institute of Physics and Technology of Materials, P.O. Box MG-7, Bucharest-Magurele, Romania b Technical University Tg.-Mures, Tg.-Mures, Romania e Institute of Technical Physics, Budapest, Hungary d INTEC, Bucharest, Romania
Received 19 March 1995;acceptedfor publication4 May 1995
Abstract
Ti nitride films were deposited by DC magnetron sputtering on HSS and Si substrates. X-ray diffraction showed the formation of single-phase g-TiN. (111) texture was found in most films, as well as stacking disorder on (111) planes, both correlated with deposition parameters, including the degree of plasma ionization. Evidence was found for plastic deformation relaxing film stress. The structure data are correlated with film microstructure and microhardness.
1. I n t r o d u c t i o n
Transition-metal nitrides and especially T i - N recently attracted considerable attention as wear-resistant coatings for high speed steel (HSS) cutting tools. A number of other potential applications emerged, ranging from diffusion barrier layers in electronics to optical coatings and transparent conducting layers in silicon solar cells. The formation of TiN x films by reactive sputtering occurs in conditions far from equilibrium, favouring the formation of metastable T i - N phases. Also, texture effects [1], dislocation-like defects [2], elastic anisotropy,
* Corresponding author. Tel.: +401 780 70 40; Fax: +401 312 22 45.
stresses in the film plane [3] strongly influence the atomic structure and microstructure of the films. Ultimately, they determine the mechanical properties (hardness, adherence, wear-resistance) which are of first-rate interest in these coatings. Structure evaluation of T i - N films by X-ray diffraction (XRD) meets with difficulties, related to the above-mentioned defects. The complexity of the situation was recently analyzed by Kuzel and coworkers [4]. This paper reports on structural properties identified by XRD in thin TiN films prepared by DC magnetron sputtering. Special attention was devoted to the effects of additional plasma ionization. The microstructure of films was studied by cross-sectional transmission electron microscopy (XTEM) and microhardness was determined for representative samples.
0169-4332/95/$09.50 © 1995 Elsevier Science B.V. All fights reserved SSDI 0169-4332(95)00134-4
296
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
2. Experimental
2.1. Film preparation The TiN layers were deposited in a laboratorybuilt installation for reactive magnetron sputtering with a rectangular Ti target ( d i m e ~ o n s 160 × 85 × 5 mm). The DC discharge was sustained by a 10 kV A source, which could be controlled (in power or current) by a microprocessor, in order to maintain a constant rate of deposition, as measured by a vibrating quartz resonator. The HSS multiphasic substrates (35 × 10 × 2 mm) were mechanically polished to a rugosity < 0.07 /xm, then cleaned in an ultrasonic bath with solvents (trichlorethylene, acetone, ethyl alcohol) and dried in N 2 flux. The single-crystal Si wafers used as substrates were etched in 20% H2SO 4, then in a solution of 5% HF in distilled water, followed by rinsing in distilled water and drying in N 2 flux. We used both HSS and Si as substrates in order to check their effect on the structure and microstructure of films. Also, samples deposited on Si could be examined by XTEM. The.substrates were placed at 60 m m distance from the target and were resistively pre-heated in vacuum (base pressure 1.3 × 10 -3 Pa). No substrate cleaning by ion sputtering was performed prior to deposition. The temperature of the substrate was measured with a miniature chromel-alumel thermocouple and was kept constant during deposition. The process parameters were controlled and adjusted by a feedback system. In a first stage, a Ti film (thickness ~ 100 nm) was deposited by sputtering in Ar atmosphere. Subsequently, the TiN layer (1-3 /xm thickness) was obtained by reactive sputtering. The acceleration voltage applied to the substrate was - 3 5 0 < Us < , 1 0 0 V. The pressure of the working gas (At + 20% N 2) was controlled dynamically by adjusting the evacuation rate and was kept constant at 0.65 Pa. The atomic ratio N / T i could be varied by controlling the sputtering power on the Ti target correlated with .the N 2 flux. However, the amount and composition of the T i - N phases are also dependent on the reactivity of the components, influenced by their ionization state. The composition of the films was studied by depth-profiling Auger analysis using the Sundgren
method [5] and was found to be uniform across film thickness. Film compositions ranged between TiN and TIN1.13. Impurity elements (O, C) were within 2 at% in most samples. For some depositions, additional plasma ionization was applied by means of an electron beam, accelerated by the ionization voltage Ua. Thereby, the ionic current at the substrate could be varied, with the sputtering power kept constant. Changing the ratio between fluxes of ionized and neutral particles bombarding the substrate could increase the yield of the TiN formation reaction. The samples investigated and their deposition parameters are shown in Table 1.
2.2. Structure investigations X-ray diffraction was performed with a 0 - 2 0 Bragg-Brentano diffractometer, equipped with a flat graphite monochromator before the NaI(T1) scintillation counter. Electronic discrimination was also used to select the C u K a radiation. Positions and widths of diffraction lines were measured in a flowchart record system, using a low angular speed of the counter ( 1 / 2 °20/min). Line positions were corrected for instrumental aberrations by means of diffraction features of the substrates (e.g. a-Fe lines of the HSS base). Full-widths at half-maximum (FWHM) were corrected for the instrumental broadening, as determined with a single crystal Si wafer. The weak TiN lines at high angles were measured with a wider-apertures system and corrected for the appropriate instrumental broadening. In all cases, corrections to FWHMs were below 6%. The combined reproducibility of preparation conditions and XRD measurements can be checked by comparing the structural data for the couples of films 2a-b, 8 a - b and 9a-b, deposited under the same conditions. The cross-sectional microstructure was investigated by XTEM, on representative samples. Microhardness measurements were performed using a Vickers microdurimeter Shimadzu, with a diamond pyramidal indentor. The values shown in Table 1 are averaged over 6 measurements taken on a 4 /zm thick film, in order to eliminate the effect of the substrate. The microhardness was calculated as HV (kg × m m -2) = 1854.4 F / d 2 ( F = force in grams, d = trace diagonal in /zm).
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
297
{111J/(2001 (areas)
3. Results
3.1. Crystalline phases
8ab~gob 8s
11S,10S
@3s
The films consisted of &TiN with cubic, NaC1type lattice and a --- 4.240 .~. No other T i - N phases could be ascertained. The XRD analysis of films on HSS was complicated by the presence in the substrate of various intermetallic phases, besides the dominant ot-Fe.
@12s
I I
@5s
3@
3.2. Preferred orientation
\x
A relatively strong texture, with planes (111) parallel to the substrate was evidenced in most samples. The degree of preferred orientation was characterized by the ratio of line areas R = (111)/(200). It turned out that R is best rationalized as a function of the substrate temperature T~ (Fig. 1). Low T~ values favour a high texture. In samples 8a,b and 9a,b the (200) line is absent, pointing to a highest degree of (111) preferred orientation. Increasing T~ above some 260°C entails a loss of preferred orientation, such that deposition on hot substrates (sample 1 with
@5
" " "~ " - r - I I
random
'00
2
300
400
Ts{* )
Fig. 1. Texture ratio (111)/(200) vs. substrate temperature Ts: (zx) Ts = 100°(2; ( O ) 250 < Ts < 270°C; ( O ) 300 < Ts < 320°C; ( [ ] ) Ts = 450°C; (zx), ( O ) samples with additional plasma ionization.
Ts = 450°C) results in R values close to that for random orientation of crystallites (R o = 0.77). Films deposited without additional ionization (samples 1,
Table 1 Preparative parameters and microhardness of TiN films Sample No.
_ Us a (V)
Ts (°C)
vs b ( A / s )
Ua e (V)
Substrate
1 2a 2b 3 4 5 6 7 8a 8b 9a 9b 10s d 1 ls 12s 3s 5s 8s
100 150 150 200 200 200 200 220 250 250 350 350 100 100 150 200 200 250
450 270 270 300 310 320 320 310 260 260 250 250 100 100 300 300 320 260
3 2.5 2.5 1.5 1.5 1.5 1.5 1.5 2.5 2.5 2.5 2.5 3 3 1.5 1.5 1.5 2.5
220 220/220/-/220 -/220 300 220 220 220/-
HSS HSS HSS HSS HSS HSS HSS HSS HSS HSS HSS HSS Single-crystal Si HSS HSS HSS HSS HSS
a Acceleration voltage. b Deposition rate. c Ionization voltage in the first and second deposition stages (15 min each). a Samples NS are deposited on Si in the same conditions as N (on HSS).
Microhardness HV20g ( k g / m m 2)
1853 2827
2012
2190
298
R. Manaila et al. / Applied Surface Science 91 (1995) 295-302
2ab, 8ab, 9ab, all with Ua = 0) as well as those with Ua = 0 in the first deposition stage (samples 6, 7) fall close to the same curve. However, additional ionization in the first (samples 4, 5) or in both deposition stages (sample 3) causes increased texture. In these cases, the R values are above those expected for their range of temperatures (300 < T~ < 320°C). It is also noticed that films deposited on Si substrates tend to have higher texture ratios than their counterparts, prepared under the same conditions on HSS (Fig. 1). These effects seem really surprising, in view of the fact that additional ionization induces structural disorder (Section 3.3). It could be explained by the increased flux of ionized particles favouring a layerby-layer growth mechanism, with the compactly packed (111) crystallographic planes parallel to the substrate. The difference between samples of types - / 2 2 0 and 2 2 0 / - points to the importance of the first deposited layers for the texture of the growing films. 3.3. Structural disorder
The degree of structural disorder in the 8-TIN lattice was estimated by the FWHMs of diffraction
B111{'281
@3
B177{°20)
(222)
÷
/7171
\ t2oo)
0
oi,
~2
2a
lines (111), (200), (311) and (222). Further lines were less reliable, due to their low intensity (caused by limited film thickness), also due to the large widths caused by disorder. Fig. 2 shows the FWHM of the (111) line vs. the acceleration voltage U~. Increasing Us (with Ua = 0) induces an improved crystallization of the TiN lattice, if taking into account films deposited between 250 and 270°C (samples 2, 8 and 9). The effect of T~
@7
a ¢~1 (222)
@6
11111
3s~/* AlOs~2
a
-go
s;n~,,x
Fig. 3. FWHM of the (111) line of 6-TIN vs. sin0/A. Samples: ([]) 2b; (zx) 3; ( Q ) 9b. Full symbol: highly oriented sample.
@ 12s
I
0',
8
t,.28
@5 (2001
~.2~
a0 (a17)
"~
• 9a, b
It
e8s L2o 0
I
700
I
I
I
200
300
~00
i
Us { VI
Fig. 2. FWHM of the (111) line of &TiN vs. acceleration voltage Ug, Legend as in Fig. 1.
~
~
sr "
Fig. 4. &TiN lattice parameter derived from different diffraction lines vs. orientation factor 3 F. Highly oriented films: ( • ) 8a; ( O ) 9b. Weakly oriented films: ([]) 2a; ( O ) 5; (zx) 6. a0: equilibrium value (for TiNt. 9 stoichiometry).
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
is also manifest. Sample 1, although deposited at a very low Us (100 V) is well crystallized, due to the high Ts used (450°C). However, the films deposited between 300 and 320°C (all of them with additional ionization) are more disordered than could be accounted for by their T~ and Us. We have to conclude to the disordering effect of the additional bombardment with charged particles caused by Us 4: 0. Films deposited on Si (with additional ionization) show a better crystallization than their counterparts on HSS. It is remarkable that not all diffraction lines are equally affected by disorder. Fig. 3 shows a selection of three samples, whose disorder increases in the sequence 9 < 2 < 3 (see Fig. 2). While the (200) and (311) lines are little sensitive to disorder, (111) and (222) are strongly affected. We conclude that structural disorder in our TiN films primarily consists in the variation of the spacing between compact planes (111), grown parallel to the substrate. At low Us and T~, this results from a low energy of the incoming clusters. In samples with additional ionization, Ar ions are presumably trapped, inducing additional layer disorder. Comparison of Fig. 1 and Fig. 2 shows that the most textured films are also the best crystallized. 3.4. Lattice parameters
The lattice parameter a of &TiN was determined from the positions of (111), (200), (311) and (222) lines for films deposited on HSS. For the former two lines, correction for instrumental aberrations was performed using the (110) line of a-Fe, situated close to their angular range (2 0 = 44 °, vs. 36 and 42 ° for the TiN lines). (311) and (222) positions were corrected by means of the a-Fe (211). This method of correction seems best suited in cases where the small number of available lines prevents the use of Nelson-Riley extrapolation procedures. In TiN films, deviations of lattice parameters are extensively used [3,4] to evaluate the compressive stress o- in the film by a sin2~ plot: ( a ¢ - a o ) / a o = A a ¢ = 2S 1o- + ( 1 / 2 ) S z o" sinZlp ( 9 = angle between the normal to substrate and reciprocal vector, S 1 and S2 = X-ray elastic constants). The dependence S 1 (3F) (with F = (h2k z + k212 h- 12h2)/(h 2 + k 2 q-/2)2 for cubic lattices), characterizing the anisotropy of polycrystailine materials,
299
can be calculated starting from the single-crystal elastic constants of TiN, using different models, of which the Hill-Neerfeld one is the most common [3]. All of these models predict a linear variation of S 1 vs. orientation factor 3F, which can be compared with the experimental trend. Fig. 4 shows the lattice parameters of 6-TIN, as derived from the four diffraction lines used, vs. the orientation factor. No linear increase a(3F) can be noticed. However, some observations are in order. The films with a low texture show positive deviations of a ( l l l ) and a(200), as expected in case of a compressive stress in the film plane. On the contrary, a(311) shows a negative deviation which is hard to explain. Highly textured samples, on the other hand, display a strong negative alteration of a ( l l l ) , although the position of line (311) is not sensitive to texture. For these highly oriented films, line (200) is absent. A large negative alteration of a(111), hardly to be explained by stress effects, was also reported in Ti(CN) films [6] and attributed to strong texture a n d / o r plastic deformation (microcracking). The same effect was found in TiN films deposited by reactive sputtering with no significant texture. Therefore, the differences seen in Fig. 4 between textured and non-textured films are to be attributed to plastic deformations of the (111) planes, relaxing the compressive stresses. In highly (111) textured films, these planes (which have a low elastic modulus) have to sustain the highest strain and easily give in. However, the decrease of a(111) below the stoichiometric value (a 0 = 4.239 for TIN0.9 [7]) is hardly accountable. It can be due, however, to deviations from stoichiometry in either direction, which were reported to decrease a 0 in bulk samples [8]. Thus, judging by the a ( l l l ) values found in highly textured samples (Fig. 4), we expectour thin films to have a bulk parameter a 0 ~ 4.22 A, which also corresponds to the a(311) values measured. This explanation agrees with depth profiling Auger analysis (Section 2.1) which reported global compositions TiNx(x > 1), i.e. an N excess over the equilibrium formula TiN0. 9. The different behaviour of lines (111) and (311) (Fig. 4) suggest that (311) is a plane of low elastic modulus, giving in (by plastic deformation) under the influence of the stress present in the films. On the other hand, the (111) planes relax only in cases of high texture, when they have to
300
R. Manaila et aL /Applied Surface Science 91 (1995) 295-302
a
b
c
d
Fig. 5. Cross-sectional micmgraphs of samples 1 ls (a), 2a (b), 7 (c) and 1 (d).
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
support the maximal elongation strain, normal to the substrate. It is to be noticed that the different behaviour of samples in Fig. 4 correlates better with their texture and not with their degree of disorder (see Fig, 2). In some samples satellites were noticed close to the (111) TiN line, as broad shoulders at ---35 ° and/or 37 ° (20). They could not be satisfactorily assigned to neither a-Ti, nor to different Ti2N phases. Also, they do not originate in the film substrate and were eliminated from the diffraction pattern by profile decomposition. In most samples, these satellites only caused an asymmetric profile of the (111) line. These effects point to a complex shape of (111) reciprocal space nodes, caused by defects (stacking faults, dislocations) a n d / o r shear stresses. 3.5. M i c r o s t r u c t u r e
The cross-sectional microstructure of selected samples deposited on Si substrates showed different aspects, dependent on the preparation parameters. Comparing microstructure with the XRD parameters (for films deposited on HSS or Si), the following conclusions emerge: 1. Strongly textured, relatively well crystallized films (10s, l l s in Table 1) deposited at low T~ (100°C) show a porous, fine-grained microstructure, with columns displaying inter-granular cavities (Fig. 5a). 2. Films with moderate texture, as revealed by XRD (2, 7) consist of a fine-grained zone, at the interface with the substrate, followed by a region of columnar formations, with a conical shape towards the film surface (Fig. 5b). 3. Sample 7 (Fig. 5c), deposited with additional ionization in the second half-stage displays an external, fine-grained zone, absent in sample 2. This region, where the columnar formations are destroyed, is caused by bombardment with energetical ions. It is to be mentioned that plasma ionization was found to increase slightly the texture of films (see Fig. 1). 4. Randomly oriented films (sample 1, see Fig. 1) show well-developed, compact columns (Fig. 5d). An apparently surprising conclusion is that welldeveloped columnar microstructure is associated with low crystallographic (111) texture. If the incoming
301
clusters get enough energy from the substrate (high T~) and no additional ionic bombardment occurs, TiN crystallites develop with random orientation, into a well-shaped columnar microstructure due to "shadowing" effects. However, at low T~ (or in the presence of strong ionic bombardment) a layer-bylayer growth mechanism is preferred, with the compact (111) planes parallel to substrate. In this case, columnar microstructure is much less developed. 3.6. M i c r o h a r d n e s s
Microhardness values (Table 1) show certain correlations with structural parameters. High defect concentrations, induced by ion bombardment (samples 12s and especially 4 in Fig. 2) seem to induce higher microhardness values. They are to be compared with films 2 and 9, deposited without additional ionization and having a better crystallization state (Fig. 2), which show lower microhardness. This defect-induced hardening is frequently encountered in metallic alloys.
4. Conclusions &TiN films with (111) texture deposited by DC magnetron sputtering show a variety of defects, including stacking disorder of (111) planes and plastic deformation on the same planes. Their structural features depend essentially on the deposition parameters: accelerating voltage Us and substrate temperature T~. Also, a strong effect is exerted by additional plasma ionization. The latter was shown to cause increasing stacking disorder of the (111) planes but also increased texture (if applied in the first stages of film growth). Thus, the physics of thin film growth under nonequilibrium conditions (incoming ionized or neutral clusters with high energy, hot substrates) is seen to be still in need of coherent physical models.
Acknowledgements The authors are deeply indebted to the Alexander von Humboldt Foundation for kindly providing computational equipment.
302
R. Manaila et a l . / Applied Surface Science 91 (1995) 295-302
References [1] A. Gokhman, Thin Solid Films 228 (1993) 229. [2] N. Zhang, W. Zhai, Y. Wang and A. Wagendristel, Vacuum 44 (1993) 51. [3] A.J. Perry, Thin Solid Films 170 (1989) 63. [4] R. Kuzel Jr., in: Proc. of the 4th Int. Syrup. TATF'94, Dresden, 1994, p. 291.
[5] J.E. Sundgren, B.O. Johansson and S.E. Karlsson, Thin Solid Films 105 (1983) 353. [6] R.Y. Fillit and A.J. Perry, Surf. Coat. Tech. 36 (1988) 647. [7] A. Christensen, J. Cryst. Growth 33 (1976) 99. [8] S. Nagakura, T. Kusuaski, F. Kaminoto and Y. Hirotsu, J. Appl. Cryst. 8 (1975) 65.
surface science ELSEVIER
Applied SurfaceScience91 (1995)295-302
Ti nitride phases in thin films deposited by DC magnetron sputtering Rodica Manaila a, ,, Domokos Biro b, Peter B. Barna c, MiNos Adamik c, Florin Zavaliche a, Stefan Craciun d, Andrei Devenyi a a Institute of Physics and Technology of Materials, P.O. Box MG-7, Bucharest-Magurele, Romania b Technical University Tg.-Mures, Tg.-Mures, Romania e Institute of Technical Physics, Budapest, Hungary d INTEC, Bucharest, Romania
Received 19 March 1995;acceptedfor publication4 May 1995
Abstract
Ti nitride films were deposited by DC magnetron sputtering on HSS and Si substrates. X-ray diffraction showed the formation of single-phase g-TiN. (111) texture was found in most films, as well as stacking disorder on (111) planes, both correlated with deposition parameters, including the degree of plasma ionization. Evidence was found for plastic deformation relaxing film stress. The structure data are correlated with film microstructure and microhardness.
1. I n t r o d u c t i o n
Transition-metal nitrides and especially T i - N recently attracted considerable attention as wear-resistant coatings for high speed steel (HSS) cutting tools. A number of other potential applications emerged, ranging from diffusion barrier layers in electronics to optical coatings and transparent conducting layers in silicon solar cells. The formation of TiN x films by reactive sputtering occurs in conditions far from equilibrium, favouring the formation of metastable T i - N phases. Also, texture effects [1], dislocation-like defects [2], elastic anisotropy,
* Corresponding author. Tel.: +401 780 70 40; Fax: +401 312 22 45.
stresses in the film plane [3] strongly influence the atomic structure and microstructure of the films. Ultimately, they determine the mechanical properties (hardness, adherence, wear-resistance) which are of first-rate interest in these coatings. Structure evaluation of T i - N films by X-ray diffraction (XRD) meets with difficulties, related to the above-mentioned defects. The complexity of the situation was recently analyzed by Kuzel and coworkers [4]. This paper reports on structural properties identified by XRD in thin TiN films prepared by DC magnetron sputtering. Special attention was devoted to the effects of additional plasma ionization. The microstructure of films was studied by cross-sectional transmission electron microscopy (XTEM) and microhardness was determined for representative samples.
0169-4332/95/$09.50 © 1995 Elsevier Science B.V. All fights reserved SSDI 0169-4332(95)00134-4
296
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
2. Experimental
2.1. Film preparation The TiN layers were deposited in a laboratorybuilt installation for reactive magnetron sputtering with a rectangular Ti target ( d i m e ~ o n s 160 × 85 × 5 mm). The DC discharge was sustained by a 10 kV A source, which could be controlled (in power or current) by a microprocessor, in order to maintain a constant rate of deposition, as measured by a vibrating quartz resonator. The HSS multiphasic substrates (35 × 10 × 2 mm) were mechanically polished to a rugosity < 0.07 /xm, then cleaned in an ultrasonic bath with solvents (trichlorethylene, acetone, ethyl alcohol) and dried in N 2 flux. The single-crystal Si wafers used as substrates were etched in 20% H2SO 4, then in a solution of 5% HF in distilled water, followed by rinsing in distilled water and drying in N 2 flux. We used both HSS and Si as substrates in order to check their effect on the structure and microstructure of films. Also, samples deposited on Si could be examined by XTEM. The.substrates were placed at 60 m m distance from the target and were resistively pre-heated in vacuum (base pressure 1.3 × 10 -3 Pa). No substrate cleaning by ion sputtering was performed prior to deposition. The temperature of the substrate was measured with a miniature chromel-alumel thermocouple and was kept constant during deposition. The process parameters were controlled and adjusted by a feedback system. In a first stage, a Ti film (thickness ~ 100 nm) was deposited by sputtering in Ar atmosphere. Subsequently, the TiN layer (1-3 /xm thickness) was obtained by reactive sputtering. The acceleration voltage applied to the substrate was - 3 5 0 < Us < , 1 0 0 V. The pressure of the working gas (At + 20% N 2) was controlled dynamically by adjusting the evacuation rate and was kept constant at 0.65 Pa. The atomic ratio N / T i could be varied by controlling the sputtering power on the Ti target correlated with .the N 2 flux. However, the amount and composition of the T i - N phases are also dependent on the reactivity of the components, influenced by their ionization state. The composition of the films was studied by depth-profiling Auger analysis using the Sundgren
method [5] and was found to be uniform across film thickness. Film compositions ranged between TiN and TIN1.13. Impurity elements (O, C) were within 2 at% in most samples. For some depositions, additional plasma ionization was applied by means of an electron beam, accelerated by the ionization voltage Ua. Thereby, the ionic current at the substrate could be varied, with the sputtering power kept constant. Changing the ratio between fluxes of ionized and neutral particles bombarding the substrate could increase the yield of the TiN formation reaction. The samples investigated and their deposition parameters are shown in Table 1.
2.2. Structure investigations X-ray diffraction was performed with a 0 - 2 0 Bragg-Brentano diffractometer, equipped with a flat graphite monochromator before the NaI(T1) scintillation counter. Electronic discrimination was also used to select the C u K a radiation. Positions and widths of diffraction lines were measured in a flowchart record system, using a low angular speed of the counter ( 1 / 2 °20/min). Line positions were corrected for instrumental aberrations by means of diffraction features of the substrates (e.g. a-Fe lines of the HSS base). Full-widths at half-maximum (FWHM) were corrected for the instrumental broadening, as determined with a single crystal Si wafer. The weak TiN lines at high angles were measured with a wider-apertures system and corrected for the appropriate instrumental broadening. In all cases, corrections to FWHMs were below 6%. The combined reproducibility of preparation conditions and XRD measurements can be checked by comparing the structural data for the couples of films 2a-b, 8 a - b and 9a-b, deposited under the same conditions. The cross-sectional microstructure was investigated by XTEM, on representative samples. Microhardness measurements were performed using a Vickers microdurimeter Shimadzu, with a diamond pyramidal indentor. The values shown in Table 1 are averaged over 6 measurements taken on a 4 /zm thick film, in order to eliminate the effect of the substrate. The microhardness was calculated as HV (kg × m m -2) = 1854.4 F / d 2 ( F = force in grams, d = trace diagonal in /zm).
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
297
{111J/(2001 (areas)
3. Results
3.1. Crystalline phases
8ab~gob 8s
11S,10S
@3s
The films consisted of &TiN with cubic, NaC1type lattice and a --- 4.240 .~. No other T i - N phases could be ascertained. The XRD analysis of films on HSS was complicated by the presence in the substrate of various intermetallic phases, besides the dominant ot-Fe.
@12s
I I
@5s
3@
3.2. Preferred orientation
\x
A relatively strong texture, with planes (111) parallel to the substrate was evidenced in most samples. The degree of preferred orientation was characterized by the ratio of line areas R = (111)/(200). It turned out that R is best rationalized as a function of the substrate temperature T~ (Fig. 1). Low T~ values favour a high texture. In samples 8a,b and 9a,b the (200) line is absent, pointing to a highest degree of (111) preferred orientation. Increasing T~ above some 260°C entails a loss of preferred orientation, such that deposition on hot substrates (sample 1 with
@5
" " "~ " - r - I I
random
'00
2
300
400
Ts{* )
Fig. 1. Texture ratio (111)/(200) vs. substrate temperature Ts: (zx) Ts = 100°(2; ( O ) 250 < Ts < 270°C; ( O ) 300 < Ts < 320°C; ( [ ] ) Ts = 450°C; (zx), ( O ) samples with additional plasma ionization.
Ts = 450°C) results in R values close to that for random orientation of crystallites (R o = 0.77). Films deposited without additional ionization (samples 1,
Table 1 Preparative parameters and microhardness of TiN films Sample No.
_ Us a (V)
Ts (°C)
vs b ( A / s )
Ua e (V)
Substrate
1 2a 2b 3 4 5 6 7 8a 8b 9a 9b 10s d 1 ls 12s 3s 5s 8s
100 150 150 200 200 200 200 220 250 250 350 350 100 100 150 200 200 250
450 270 270 300 310 320 320 310 260 260 250 250 100 100 300 300 320 260
3 2.5 2.5 1.5 1.5 1.5 1.5 1.5 2.5 2.5 2.5 2.5 3 3 1.5 1.5 1.5 2.5
220 220/220/-/220 -/220 300 220 220 220/-
HSS HSS HSS HSS HSS HSS HSS HSS HSS HSS HSS HSS Single-crystal Si HSS HSS HSS HSS HSS
a Acceleration voltage. b Deposition rate. c Ionization voltage in the first and second deposition stages (15 min each). a Samples NS are deposited on Si in the same conditions as N (on HSS).
Microhardness HV20g ( k g / m m 2)
1853 2827
2012
2190
298
R. Manaila et al. / Applied Surface Science 91 (1995) 295-302
2ab, 8ab, 9ab, all with Ua = 0) as well as those with Ua = 0 in the first deposition stage (samples 6, 7) fall close to the same curve. However, additional ionization in the first (samples 4, 5) or in both deposition stages (sample 3) causes increased texture. In these cases, the R values are above those expected for their range of temperatures (300 < T~ < 320°C). It is also noticed that films deposited on Si substrates tend to have higher texture ratios than their counterparts, prepared under the same conditions on HSS (Fig. 1). These effects seem really surprising, in view of the fact that additional ionization induces structural disorder (Section 3.3). It could be explained by the increased flux of ionized particles favouring a layerby-layer growth mechanism, with the compactly packed (111) crystallographic planes parallel to the substrate. The difference between samples of types - / 2 2 0 and 2 2 0 / - points to the importance of the first deposited layers for the texture of the growing films. 3.3. Structural disorder
The degree of structural disorder in the 8-TIN lattice was estimated by the FWHMs of diffraction
B111{'281
@3
B177{°20)
(222)
÷
/7171
\ t2oo)
0
oi,
~2
2a
lines (111), (200), (311) and (222). Further lines were less reliable, due to their low intensity (caused by limited film thickness), also due to the large widths caused by disorder. Fig. 2 shows the FWHM of the (111) line vs. the acceleration voltage U~. Increasing Us (with Ua = 0) induces an improved crystallization of the TiN lattice, if taking into account films deposited between 250 and 270°C (samples 2, 8 and 9). The effect of T~
@7
a ¢~1 (222)
@6
11111
3s~/* AlOs~2
a
-go
s;n~,,x
Fig. 3. FWHM of the (111) line of 6-TIN vs. sin0/A. Samples: ([]) 2b; (zx) 3; ( Q ) 9b. Full symbol: highly oriented sample.
@ 12s
I
0',
8
t,.28
@5 (2001
~.2~
a0 (a17)
"~
• 9a, b
It
e8s L2o 0
I
700
I
I
I
200
300
~00
i
Us { VI
Fig. 2. FWHM of the (111) line of &TiN vs. acceleration voltage Ug, Legend as in Fig. 1.
~
~
sr "
Fig. 4. &TiN lattice parameter derived from different diffraction lines vs. orientation factor 3 F. Highly oriented films: ( • ) 8a; ( O ) 9b. Weakly oriented films: ([]) 2a; ( O ) 5; (zx) 6. a0: equilibrium value (for TiNt. 9 stoichiometry).
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
is also manifest. Sample 1, although deposited at a very low Us (100 V) is well crystallized, due to the high Ts used (450°C). However, the films deposited between 300 and 320°C (all of them with additional ionization) are more disordered than could be accounted for by their T~ and Us. We have to conclude to the disordering effect of the additional bombardment with charged particles caused by Us 4: 0. Films deposited on Si (with additional ionization) show a better crystallization than their counterparts on HSS. It is remarkable that not all diffraction lines are equally affected by disorder. Fig. 3 shows a selection of three samples, whose disorder increases in the sequence 9 < 2 < 3 (see Fig. 2). While the (200) and (311) lines are little sensitive to disorder, (111) and (222) are strongly affected. We conclude that structural disorder in our TiN films primarily consists in the variation of the spacing between compact planes (111), grown parallel to the substrate. At low Us and T~, this results from a low energy of the incoming clusters. In samples with additional ionization, Ar ions are presumably trapped, inducing additional layer disorder. Comparison of Fig. 1 and Fig. 2 shows that the most textured films are also the best crystallized. 3.4. Lattice parameters
The lattice parameter a of &TiN was determined from the positions of (111), (200), (311) and (222) lines for films deposited on HSS. For the former two lines, correction for instrumental aberrations was performed using the (110) line of a-Fe, situated close to their angular range (2 0 = 44 °, vs. 36 and 42 ° for the TiN lines). (311) and (222) positions were corrected by means of the a-Fe (211). This method of correction seems best suited in cases where the small number of available lines prevents the use of Nelson-Riley extrapolation procedures. In TiN films, deviations of lattice parameters are extensively used [3,4] to evaluate the compressive stress o- in the film by a sin2~ plot: ( a ¢ - a o ) / a o = A a ¢ = 2S 1o- + ( 1 / 2 ) S z o" sinZlp ( 9 = angle between the normal to substrate and reciprocal vector, S 1 and S2 = X-ray elastic constants). The dependence S 1 (3F) (with F = (h2k z + k212 h- 12h2)/(h 2 + k 2 q-/2)2 for cubic lattices), characterizing the anisotropy of polycrystailine materials,
299
can be calculated starting from the single-crystal elastic constants of TiN, using different models, of which the Hill-Neerfeld one is the most common [3]. All of these models predict a linear variation of S 1 vs. orientation factor 3F, which can be compared with the experimental trend. Fig. 4 shows the lattice parameters of 6-TIN, as derived from the four diffraction lines used, vs. the orientation factor. No linear increase a(3F) can be noticed. However, some observations are in order. The films with a low texture show positive deviations of a ( l l l ) and a(200), as expected in case of a compressive stress in the film plane. On the contrary, a(311) shows a negative deviation which is hard to explain. Highly textured samples, on the other hand, display a strong negative alteration of a ( l l l ) , although the position of line (311) is not sensitive to texture. For these highly oriented films, line (200) is absent. A large negative alteration of a(111), hardly to be explained by stress effects, was also reported in Ti(CN) films [6] and attributed to strong texture a n d / o r plastic deformation (microcracking). The same effect was found in TiN films deposited by reactive sputtering with no significant texture. Therefore, the differences seen in Fig. 4 between textured and non-textured films are to be attributed to plastic deformations of the (111) planes, relaxing the compressive stresses. In highly (111) textured films, these planes (which have a low elastic modulus) have to sustain the highest strain and easily give in. However, the decrease of a(111) below the stoichiometric value (a 0 = 4.239 for TIN0.9 [7]) is hardly accountable. It can be due, however, to deviations from stoichiometry in either direction, which were reported to decrease a 0 in bulk samples [8]. Thus, judging by the a ( l l l ) values found in highly textured samples (Fig. 4), we expectour thin films to have a bulk parameter a 0 ~ 4.22 A, which also corresponds to the a(311) values measured. This explanation agrees with depth profiling Auger analysis (Section 2.1) which reported global compositions TiNx(x > 1), i.e. an N excess over the equilibrium formula TiN0. 9. The different behaviour of lines (111) and (311) (Fig. 4) suggest that (311) is a plane of low elastic modulus, giving in (by plastic deformation) under the influence of the stress present in the films. On the other hand, the (111) planes relax only in cases of high texture, when they have to
300
R. Manaila et aL /Applied Surface Science 91 (1995) 295-302
a
b
c
d
Fig. 5. Cross-sectional micmgraphs of samples 1 ls (a), 2a (b), 7 (c) and 1 (d).
R. Manaila et al. /Applied Surface Science 91 (1995) 295-302
support the maximal elongation strain, normal to the substrate. It is to be noticed that the different behaviour of samples in Fig. 4 correlates better with their texture and not with their degree of disorder (see Fig, 2). In some samples satellites were noticed close to the (111) TiN line, as broad shoulders at ---35 ° and/or 37 ° (20). They could not be satisfactorily assigned to neither a-Ti, nor to different Ti2N phases. Also, they do not originate in the film substrate and were eliminated from the diffraction pattern by profile decomposition. In most samples, these satellites only caused an asymmetric profile of the (111) line. These effects point to a complex shape of (111) reciprocal space nodes, caused by defects (stacking faults, dislocations) a n d / o r shear stresses. 3.5. M i c r o s t r u c t u r e
The cross-sectional microstructure of selected samples deposited on Si substrates showed different aspects, dependent on the preparation parameters. Comparing microstructure with the XRD parameters (for films deposited on HSS or Si), the following conclusions emerge: 1. Strongly textured, relatively well crystallized films (10s, l l s in Table 1) deposited at low T~ (100°C) show a porous, fine-grained microstructure, with columns displaying inter-granular cavities (Fig. 5a). 2. Films with moderate texture, as revealed by XRD (2, 7) consist of a fine-grained zone, at the interface with the substrate, followed by a region of columnar formations, with a conical shape towards the film surface (Fig. 5b). 3. Sample 7 (Fig. 5c), deposited with additional ionization in the second half-stage displays an external, fine-grained zone, absent in sample 2. This region, where the columnar formations are destroyed, is caused by bombardment with energetical ions. It is to be mentioned that plasma ionization was found to increase slightly the texture of films (see Fig. 1). 4. Randomly oriented films (sample 1, see Fig. 1) show well-developed, compact columns (Fig. 5d). An apparently surprising conclusion is that welldeveloped columnar microstructure is associated with low crystallographic (111) texture. If the incoming
301
clusters get enough energy from the substrate (high T~) and no additional ionic bombardment occurs, TiN crystallites develop with random orientation, into a well-shaped columnar microstructure due to "shadowing" effects. However, at low T~ (or in the presence of strong ionic bombardment) a layer-bylayer growth mechanism is preferred, with the compact (111) planes parallel to substrate. In this case, columnar microstructure is much less developed. 3.6. M i c r o h a r d n e s s
Microhardness values (Table 1) show certain correlations with structural parameters. High defect concentrations, induced by ion bombardment (samples 12s and especially 4 in Fig. 2) seem to induce higher microhardness values. They are to be compared with films 2 and 9, deposited without additional ionization and having a better crystallization state (Fig. 2), which show lower microhardness. This defect-induced hardening is frequently encountered in metallic alloys.
4. Conclusions &TiN films with (111) texture deposited by DC magnetron sputtering show a variety of defects, including stacking disorder of (111) planes and plastic deformation on the same planes. Their structural features depend essentially on the deposition parameters: accelerating voltage Us and substrate temperature T~. Also, a strong effect is exerted by additional plasma ionization. The latter was shown to cause increasing stacking disorder of the (111) planes but also increased texture (if applied in the first stages of film growth). Thus, the physics of thin film growth under nonequilibrium conditions (incoming ionized or neutral clusters with high energy, hot substrates) is seen to be still in need of coherent physical models.
Acknowledgements The authors are deeply indebted to the Alexander von Humboldt Foundation for kindly providing computational equipment.
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References [1] A. Gokhman, Thin Solid Films 228 (1993) 229. [2] N. Zhang, W. Zhai, Y. Wang and A. Wagendristel, Vacuum 44 (1993) 51. [3] A.J. Perry, Thin Solid Films 170 (1989) 63. [4] R. Kuzel Jr., in: Proc. of the 4th Int. Syrup. TATF'94, Dresden, 1994, p. 291.
[5] J.E. Sundgren, B.O. Johansson and S.E. Karlsson, Thin Solid Films 105 (1983) 353. [6] R.Y. Fillit and A.J. Perry, Surf. Coat. Tech. 36 (1988) 647. [7] A. Christensen, J. Cryst. Growth 33 (1976) 99. [8] S. Nagakura, T. Kusuaski, F. Kaminoto and Y. Hirotsu, J. Appl. Cryst. 8 (1975) 65.