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Structure of deposited and annealed TiB2 layers
Surface and Coatings Technology 97 (1997) 313–321
Structure of deposited and annealed TiB layers 2 Renate Wiedemann *, Heinrich Oettel, Marko Jerenz Freiberg University of Mining and Technology, Institute of Physical Metallurgy, Gustav-Zeuner-Strasse 5, D-09596 Freiberg, Germany
Abstract TiB coatings were prepared on steel and silicon substrates by d.c. magnetron sputtering at different bias voltages. After 2 deposition part of the coatings was vacuum annealed at 400 °C and 800 °C. Transmission elctron microscopic (TEM ) investigations on cross-section-prepared specimens showed that the layers were nanocrystalline with an average diameter of columnar grains between 50 and 20 nm depending on bias voltage. The chemical composition of coatings was homogeneous within the layers and independent of bias. A non-stoichiometric B:Ti ratio was detectable and a high amount of Ar incorporation occurred. X-Ray phase analysis showed that the coatings consisted mainly of hexagonal TiB phase with strong (001) fibre texture. Moreover, high 2 compressive stresses were measured which could be attributed to Ar incorporation. The microhardness, critical load and failure mode were influenced by high compressive residual stresses. After annealing at 400 °C the residual stresses were relaxed and the critical load was independent of bias voltage. After annealing at 800 °C an upwelling of the surface was observed connected with crack formations and the occurrence of three new phases. © 1997 Elsevier Science S.A. Keywords: Thin hard coating; Microstructure; Chemical composition; Annealing; Mechanical properties
1. Introduction
2. Experimental details
TiB is a hexagonal compound with metallic chemical 2 bonding character. In comparison to the well-known TiN it has a lower expansion coefficient and a better adhesion to metallic substrates but also a higher interaction tendency with other materials. Moreover, it is characterized by a high melting point (3225 °C ) and high hardness value (~3000 HV ) and shows good electrical conductivity. However, the properties are anisotropic. They can differ up to 30% depending on orientation [1]. TiB has only a small homogeneity 2 range in the binary system, i.e. the B:Ti ratio is nearly constant under equilibrium conditions. However, the properties of coatings can differ from those of the bulk material. Some publications have reported on the microstructure and properties of TiB produced under various 2 conditions on different substrates. However, the results are insufficient and sometimes contradictory. The aim of our work is to investigate the microstructure and properties of sputtered TiB coatings in dependence on 2 various bias voltages and their annealing behaviour to test the application limits.
TiB layers were deposited on steel substrates (heat2 treatable steel 43 CrMo 4) and oxidized Si(001) wafers by d.c. magnetron sputtering using a commercial magnetron sputter equipment (Balzers PLS 500). A 97.5% TiB target was used and the sputter conditions were 2 adjusted on 600 W power control and 0.5 Pa Ar gas pressure. The bias voltage was changed in the range of 0 to −200 V. A r.f. sputter cleaning was carried out before deposition. The surfaces of all steel samples were ground and polished. Some Si substrates were pretreated with an Al interlayer to improve the adhesion. A quantitative depth profile of the chemical composition was determined using a GDOES (LECO type 750/SDP). The Ar content was measured by electron beam microanalysis. The morphology was studied by scanning electron microscopy (SEM ) on cracked samples and by transmission electron microscopy ( TEM ) on cross-section-prepared layers deposited on silicon substrates. The diameter of columnar grains was determined semiquantitatively by an image analysing system. The residual stresses in the coatings were calculated
* Corresponding author. 0257-8972/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 02 5 7 -8 9 7 2 ( 9 7 ) 0 0 20 4 - 1
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from the bending of samples using the STONEY formula. The bending was measured by deflection of two parallel laser-beams. Silicon wafers without an interlayer were used as substrates. The X-ray diffraction ( XRD) spectra of the coatings were recorded with a Seifert-FPM XRD 7 powder diffractometer using Cu–Ka radiation and symmetric (standard Bragg-Brentano diffraction geometry) and asymmetric beam path (Seemann-Bohlin diffraction geometry, incidence angles 4° and 8°). Quantitative pole figures were obtained by texture goniometer measurement with Cu–Ka radiation on (001) and (102) peaks. The microhardness testing was carried out by Vickers hardness measurement under load (Shimadzu) with different loads (5, 10, 25, 50, 100 p). The adhesion was estimated by a scratch-test with Rockwell-indenter and continuous increasing load up to 100 N (Revetest from CSEM ). The critical load was attributed to the normal force at which first delaminations occur (SEM, 500:1).
3. Results and discussion 3.1. Morphology Fig. 1 shows a representative SEM fractogram of the TiB –coating morphology. In contrast to other authors 2 [2–4], no columnar structure could be observed. All samples had the same structure as amorphous materials, independent of the bias voltage. However, TEM investigations on cross-section-prepared samples demonstrated that these coatings consist of very fine columnar grains (Fig. 2). The average column diameter increased from the interface to the surface of layers. It amounted on the interface to <5 nm and on the surface to approx. 40–50 nm. Smaller values were measured at higher bias voltages. It was impossible to estimate the column size
Fig. 1. Representative SEM fracture cross-section of TiB coating. 2
Fig. 2. TEM imaging of cross-section-prepared TiB coating. 2
because grain boundaries could not be detected in the growth direction of the grains. The layer thickness was measured on fractured samples and amounted about 2.5 mm on all samples. This means that no significant influence of bias voltage could be observed. The deposition rate calculated from layer thickness was 0.7 nm s−1. 3.2. Chemical composition The GDOES concentration–depth profiles of three layers deposited at 0, −80 and −200 V bias voltages are shown in Fig. 3. A homogeneous contribution of
Fig. 3. GDOES concentration–depth profile of three TiB coatings. 2
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titanium and boron could be observed within the layers. In contrast with values in the literature [3,5] the contents of titanium and boron were independent of bias voltage. Moreover, the content of boron in all samples was very high, estimated at 70 at.%. The B:Ti ratio of films was 2.6 and accordingly higher than the stoichiometric ratio. The high boron content could not be attributed to measuring or standardization errors, and it has also been detected by electron beam microanalysis measurements. Comparable B:Ti ratios have been determined previously [2–4]. Electron beam microanalysis showed a remarkable Ar incorporation in all samples. The Ar content increased with increasing bias voltage from 0.18 at.% at 0 V bias to 2.4 at.% at 200 V. It increased rapidly above −50 V bias, i.e. with higher energy of argon ions.
3.3. Microstructure Figs. 4 and 5 show representative X-ray diffraction patterns and the pertinent quantitative pole figures of three coatings deposited with various bias voltages. All layers produced with low bias consisted of a hexagonal TiB phase with a strong (001) fibre texture and a small 2 quantity of the orthorhombic Ti B phase. With increas3 4 ing bias the integral intensities of TiB peaks decreased 2 and the integral breadth increased ( Table 1), because of the smaller grain size. At −200 V bias the diffraction pattern showed only wide TiB peaks with very low 2 intensities, and the pole figures demonstrated that the layer crystals were more randomly oriented. Detection of Ti B was impossible in these coatings. 3 4 A crystalline boron phase could not be detected in any coating by either X-ray or TEM investigations. High compressive residual stresses occurred in all layers. Stresses of −1.4 GPa were measured also in coatings deposited without bias. With increasing bias the residual stress increased to −4.5 GPa at 150 V. The amount of stress generated was dependent on the level of argon content but not on the amount of boron (Fig. 6). Therefore, Ar incorporation can be regarded as the main reason for stress formation. Unfortunately, the adhesion of coatings deposited with higher bias voltages on silicon substrates was not sufficient. Partial delaminations took place and, therefore, the measurement of residual stresses by bending was impossible. Strong texture, high compressive stresses and low intensities cause great problems for the estimation of the correct lattice constants. Nevertheless, measured values can be compared with one another. The following lattice parameters of TiB phase were calculated: 2 0 V bias: a=0.3051±0.0011 nm, c=0.3193±0.0012 nm; −200 V bias: a=0.3044±0.0010 nm, c=0.3224± 0.0011 nm.
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The errors were relatively great as a consequence of difficulties mentioned above. The measured values were not stress corrected and, therefore, they are not comparable with bulk material directly (bulk: a=0.303034 nm, c=0.322953 nm [6 ]). However, the lattice parameter a decreased and the parameter c increased with rising bias corresponding to higher residual stresses, meaning that the hexagonal lattice was stretched in (001) direction. It is not clear where argon and the excess boron are placed. In the hexagonal TiB lattice, incorporation of 2 interstitial atoms is only possible in tetrahedral holes. However, the misfit of hole and argon radii is about 60% and Ar incorporation is unsuccessful. The incorporation of boron in tetrahedral holes can also be excluded. Substitution of Ti atoms by argon or boron atoms is also improbable owing to the large difference in the radii (r =0.14 nm, r =0.10 nm, r =0.09 nm) and Ti Ar B would produce tensile stresses. In the TiN lattice Ar atoms substitute N atoms under generation of compressive stresses [7]. The same result should be possible in TiB when Ar atoms occupy boron 2 sites, but this reflection does not agree with the measured high boron content.
3.4. Mechanical properties Direct measurement of conventional Vickers hardness is impossible because the layers are very brittle and many cracks are formed in and around all indentations. Therefore, the diagonals of indentations cannot be determined correctly. For that reason, hardness testing under load is used where the indentation depth of a Vickers indenter is measured. Commonly, in thin layers the direct determination of layer hardness is impossible because the indentation depth is larger than the layer thickness. It is possible to calculate layer hardness without the influence of substrate hardness, using a method described in Ref. [8]. In this case, the measured compound hardness values estimated at different loads are expressed as a function of the ratio of the layer thickness to the indentation depth. The layer hardness can be calculated from the slope of the corresponding curve. Using this method the following layer hardness values were determined: 0 V bias: HU=31 500 N mm−2; −200 V bias: HU= 34 300 N mm−2. The results show that the high compressive residual stresses increase the hardness of TiB layers. Similar 2 results on the influence of residual stresses have been presented previously [9,10]. However, the measured hardness values of TiB fluctuate between 3000 and 2 6700 HV. This is caused by the difficulties of thin-layer hardness measurements which are described above and more in detail in Ref. [8].
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Fig. 4. X-Ray diffraction patterns.
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Fig. 5. Quantitative pole figures.
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Table 1 Integral intensities (I ) and integral breadth of TiB layers deposited 2 with various bias voltages Bias voltage (V )
I (deg s−1)
B (°2H)
0 −50 −200
1106 286 8
0.32 0.39 0.91
stresses generated as a sequence of soft substrate material displacements during penetration of the stylus. Higher compressive stresses prevent cohesive cracking under the stylus and favour spalling or buckling in front of the indenter. Fig. 7 gives typical examples of different failure modes. Our results are in contrast to previous measurements [3], where the critical load decreased with increasing bias. However, for further discussion more information on failure modes and coating properties is required.
3.5. Annealing
Fig. 6. Residual stresses and argon content in dependence on bias voltage.
The compressive residual stresses also improve the adhesion of layers and influence the failure modes during scratch testing. The critical load of adhesive failure increases with rising bias from 4 N at 0 V to 19 N at −200 V. In general, the measured critical load values are very low, because of the soft substrate. The substrate hardness amounts only 380 HV 0.1. Other authors [4] have measured 50–90 N for TiB layers on cemented 2 carbide. The failure of coatings deposited with lower compressive stresses (0 V bias) takes place with long cracks starting in the middle and small shell-shaped delaminations at the edges of scratch tracks. Samples with higher compressive stresses exhibit only small cracks at the track edges using the same normal force. These layers fail by upwelling and spalling in front of the stylus at higher forces. The first type of cracks and small delaminations are caused by tensile
After annealing at 400 °C the chemical composition and the microstructure (phases and texture) are not modified and the critical load of adhesive failure is independent of the bias voltage. The values of all investigated samples amount to only 3–4 N and, thus, they are placed in the range of 0 V bias samples before annealing. It is supposed that in this state the compressive residual stresses produced during deposition should be relaxed. Therefore, these stresses do not improve the adhesion. Nevertheless, the failure mode is the same as before annealing, which may indicate that the failure mode is also influenced by morphology (smaller grain size with increasing bias). During heating to 400 °C tensile stresses up to 1.3 GPa are produced, which are caused by the difference between the thermal expansion coefficients of iron and TiB 2 (a =12 · 10−6 K−1, a =7.8 10−6 K−1), but cracks Fe TiB2 could not be observed in the layers. Accordingly, the layers are able to compensate these stresses elastically. After 800 °C annealing the morphology of all coatings is modified. The surface is divided by many cracks because tensile stresses grow to 2.6 GPa during heating. The crack pattern is generated periodically. Furthermore, an upwelling of the surface between the cracks can be observed. The cross-sections show three layers: two (I, II ) in the former coating and one (III ) with large and lengthy grains and large pores under the former interface between layer and substrate (Fig. 8). The thickness of layer I and II is in the same range as the TiB layer before annealing. Layer III is up to 2 6 mm thick. The GDOES quantitative concentration–depth profiles exhibit a new distribution of B, C and Fe, i.e. distant diffusion occurs (Fig. 9). Boron diffuses up to 10 mm into the steel substrate. The profiles can also be divided in three parts: layer I, on the surface with high boron content; layer II, below, with a maximum of boron and titanium and a minimum of boron content; and layer III, in the former steel substrate with a second maximum of boron content. The X-ray phase diagram of symmetric X-ray diffraction also contains TiB and three new phases, TiC, FeTi 2
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Fig. 7. Failure modes during scratch testing (above, low compressive stresses; below, high compressive stresses).
and Fe B. The X-ray phase analyses with glancing 2 incidence at different angles prove that the phase Fe B 2 only occurs under the former interface. From the concentration–depth profiles and the X-ray phase analyses it can be concluded that the first layer with the highest boron content should consist mainly of TiB . It can be assumed that TiC occurs in the second 2 layer where a minimum of boron and a maximum of carbon content exist. Beside this, the formation of FeTi is possible. In the
upper range of steel substrate the ferrite is completely transformed in Fe B, which is connected with new grain 2 formation and growth. The phase Fe C is modified in 3 Fe B and free carbon which diffuses in the former 2 Ti B layer and combines with free titanium. Therefore, 2 from surface to substrate the following phases could be formed: layer I, TiB ; layer II, TiC, FeTi, also possibly 2 TiB ; layer III, Fe B. 2 2 The generation of the new phases is a consequence of the free enthalpy reduction. The following chemical
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Fig. 8. SEM imaging of fracture cross-section after annealing (above, SE contrast; below, BSE contrast).
Fig. 9. GDOES concentration–depth profile after annealing (I–III, layer thickness measured at fractured samples, see Fig. 8).
reactions take place: 3TiB +6Fe+Fe CTiB +TiC+TiFe+4Fe B 2 3 2 2 S°G=−302 kJ K mol−1 S°G=−384 kJ K mol−1 D°G=−82 kJ K mol−1. The layers are also very brittle and not suitable for application.
4. Conclusions In this investigation the structure of TiB layers 2 deposited with magnetron sputtering at −200 V bias
was nanocrystalline (grain diameters 5–40 nm) and not amorphous. The boron content was higher than in the stoichiometric one, but no crystalline boron phases could be detected by XRD or TEM. The occurrence of the Ti B phase 3 4 could not explain the high boron content. It should be excluded that boron atoms are incorporated in interstitial holes or on regular Ti atom sites. The authors of Ref. [4] supposed that the excess B atoms are incorporated in the grain boundaries of nanocrystals. The incorporation of Ar atoms in the hexagonal TiB lattice is regarded as the main reason for the 2 generation of high compressive stresses in sputtered TiB coatings. It is unclear on which sites in the lattice 2 the argon atoms are placed. Higher compressive stresses increase the adhesion and the hardness of TiB layers. The value of critical load 2 should be improved by harder substrates. TiB layers react with steel substrates at T>400 °C. 2 A diffusion barrier layer should be used for applications at higher temperatures, but crack formation caused by various thermal coefficients will also occur. These features limit the application temperature of these layers.
References [1] H. Hollek, Material selection for hard coatings, J. Vac. Sci. Technol. A4 (6) (1986) 2661–2669. [2] W. Gissler, Structure and properties of Ti–B–N coatings, Surf. Coat. Technol. 68/69 (1994) 556–563. [3] R. Lohmann, et al., Analysis of r.f.-sputtered TiB hard coatings 2
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[4]
[5] [6 ] [7]
by means of X-ray diffractometry and Auger electron spectroscopy, Mat. Sci. Engng A139 (1991) 259–263. ¨ C. Mitterer, M. Rauter, P. Rodhammer, Sputter deposition of ultrahard coatings within the system Ti–B–C–N, Surf. Coat. Technol. 41 (1990) 351–363. T. Shikama, et al., Deposition of TiB films by a co-sputtering 2 method, Thin Solid Films 156 (1988) 287–293. Powder Diffraction File PDF-2 Database Sets 1–42, International Centre for Diffraction Data (JCPDS), 1992. R. Wiedemann, H. Oettel, Macroscopic and microscopic stresses in nitride hard coatings, in: V. Hauk et al., Residual Stresses; DGM Informationsges, 1993, 673–681.
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¨ [8] A. Kirsten, R. Wiedemann, H. Oettel, Ermittlung der Schichtharte ¨ ¨ aus gemesenen Verbundhartewerten, Harterei-Techn. Mitt. 48 (1993) 324–329. [9] W. Herr, B. Matthes, E. Broszeit, K.H. Kloos, Fundamental properties and wear resistance of r.f. sputtered TiB and Ti(B,N ) 2 coatings, Mat. Sci. Engng A140 (1991) 616–624. [10] O. Knotek, F. Jungblut, R. Breitenbach, Magnetron-sputtered superhard coatings within the system Ti–B–C–N, Vacuum 41 (1990) 2184–2186.
Structure of deposited and annealed TiB layers 2 Renate Wiedemann *, Heinrich Oettel, Marko Jerenz Freiberg University of Mining and Technology, Institute of Physical Metallurgy, Gustav-Zeuner-Strasse 5, D-09596 Freiberg, Germany
Abstract TiB coatings were prepared on steel and silicon substrates by d.c. magnetron sputtering at different bias voltages. After 2 deposition part of the coatings was vacuum annealed at 400 °C and 800 °C. Transmission elctron microscopic (TEM ) investigations on cross-section-prepared specimens showed that the layers were nanocrystalline with an average diameter of columnar grains between 50 and 20 nm depending on bias voltage. The chemical composition of coatings was homogeneous within the layers and independent of bias. A non-stoichiometric B:Ti ratio was detectable and a high amount of Ar incorporation occurred. X-Ray phase analysis showed that the coatings consisted mainly of hexagonal TiB phase with strong (001) fibre texture. Moreover, high 2 compressive stresses were measured which could be attributed to Ar incorporation. The microhardness, critical load and failure mode were influenced by high compressive residual stresses. After annealing at 400 °C the residual stresses were relaxed and the critical load was independent of bias voltage. After annealing at 800 °C an upwelling of the surface was observed connected with crack formations and the occurrence of three new phases. © 1997 Elsevier Science S.A. Keywords: Thin hard coating; Microstructure; Chemical composition; Annealing; Mechanical properties
1. Introduction
2. Experimental details
TiB is a hexagonal compound with metallic chemical 2 bonding character. In comparison to the well-known TiN it has a lower expansion coefficient and a better adhesion to metallic substrates but also a higher interaction tendency with other materials. Moreover, it is characterized by a high melting point (3225 °C ) and high hardness value (~3000 HV ) and shows good electrical conductivity. However, the properties are anisotropic. They can differ up to 30% depending on orientation [1]. TiB has only a small homogeneity 2 range in the binary system, i.e. the B:Ti ratio is nearly constant under equilibrium conditions. However, the properties of coatings can differ from those of the bulk material. Some publications have reported on the microstructure and properties of TiB produced under various 2 conditions on different substrates. However, the results are insufficient and sometimes contradictory. The aim of our work is to investigate the microstructure and properties of sputtered TiB coatings in dependence on 2 various bias voltages and their annealing behaviour to test the application limits.
TiB layers were deposited on steel substrates (heat2 treatable steel 43 CrMo 4) and oxidized Si(001) wafers by d.c. magnetron sputtering using a commercial magnetron sputter equipment (Balzers PLS 500). A 97.5% TiB target was used and the sputter conditions were 2 adjusted on 600 W power control and 0.5 Pa Ar gas pressure. The bias voltage was changed in the range of 0 to −200 V. A r.f. sputter cleaning was carried out before deposition. The surfaces of all steel samples were ground and polished. Some Si substrates were pretreated with an Al interlayer to improve the adhesion. A quantitative depth profile of the chemical composition was determined using a GDOES (LECO type 750/SDP). The Ar content was measured by electron beam microanalysis. The morphology was studied by scanning electron microscopy (SEM ) on cracked samples and by transmission electron microscopy ( TEM ) on cross-section-prepared layers deposited on silicon substrates. The diameter of columnar grains was determined semiquantitatively by an image analysing system. The residual stresses in the coatings were calculated
* Corresponding author. 0257-8972/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 02 5 7 -8 9 7 2 ( 9 7 ) 0 0 20 4 - 1
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from the bending of samples using the STONEY formula. The bending was measured by deflection of two parallel laser-beams. Silicon wafers without an interlayer were used as substrates. The X-ray diffraction ( XRD) spectra of the coatings were recorded with a Seifert-FPM XRD 7 powder diffractometer using Cu–Ka radiation and symmetric (standard Bragg-Brentano diffraction geometry) and asymmetric beam path (Seemann-Bohlin diffraction geometry, incidence angles 4° and 8°). Quantitative pole figures were obtained by texture goniometer measurement with Cu–Ka radiation on (001) and (102) peaks. The microhardness testing was carried out by Vickers hardness measurement under load (Shimadzu) with different loads (5, 10, 25, 50, 100 p). The adhesion was estimated by a scratch-test with Rockwell-indenter and continuous increasing load up to 100 N (Revetest from CSEM ). The critical load was attributed to the normal force at which first delaminations occur (SEM, 500:1).
3. Results and discussion 3.1. Morphology Fig. 1 shows a representative SEM fractogram of the TiB –coating morphology. In contrast to other authors 2 [2–4], no columnar structure could be observed. All samples had the same structure as amorphous materials, independent of the bias voltage. However, TEM investigations on cross-section-prepared samples demonstrated that these coatings consist of very fine columnar grains (Fig. 2). The average column diameter increased from the interface to the surface of layers. It amounted on the interface to <5 nm and on the surface to approx. 40–50 nm. Smaller values were measured at higher bias voltages. It was impossible to estimate the column size
Fig. 1. Representative SEM fracture cross-section of TiB coating. 2
Fig. 2. TEM imaging of cross-section-prepared TiB coating. 2
because grain boundaries could not be detected in the growth direction of the grains. The layer thickness was measured on fractured samples and amounted about 2.5 mm on all samples. This means that no significant influence of bias voltage could be observed. The deposition rate calculated from layer thickness was 0.7 nm s−1. 3.2. Chemical composition The GDOES concentration–depth profiles of three layers deposited at 0, −80 and −200 V bias voltages are shown in Fig. 3. A homogeneous contribution of
Fig. 3. GDOES concentration–depth profile of three TiB coatings. 2
R. Wiedemann et al. / Surface and Coatings Technology 97 (1997) 313–321
titanium and boron could be observed within the layers. In contrast with values in the literature [3,5] the contents of titanium and boron were independent of bias voltage. Moreover, the content of boron in all samples was very high, estimated at 70 at.%. The B:Ti ratio of films was 2.6 and accordingly higher than the stoichiometric ratio. The high boron content could not be attributed to measuring or standardization errors, and it has also been detected by electron beam microanalysis measurements. Comparable B:Ti ratios have been determined previously [2–4]. Electron beam microanalysis showed a remarkable Ar incorporation in all samples. The Ar content increased with increasing bias voltage from 0.18 at.% at 0 V bias to 2.4 at.% at 200 V. It increased rapidly above −50 V bias, i.e. with higher energy of argon ions.
3.3. Microstructure Figs. 4 and 5 show representative X-ray diffraction patterns and the pertinent quantitative pole figures of three coatings deposited with various bias voltages. All layers produced with low bias consisted of a hexagonal TiB phase with a strong (001) fibre texture and a small 2 quantity of the orthorhombic Ti B phase. With increas3 4 ing bias the integral intensities of TiB peaks decreased 2 and the integral breadth increased ( Table 1), because of the smaller grain size. At −200 V bias the diffraction pattern showed only wide TiB peaks with very low 2 intensities, and the pole figures demonstrated that the layer crystals were more randomly oriented. Detection of Ti B was impossible in these coatings. 3 4 A crystalline boron phase could not be detected in any coating by either X-ray or TEM investigations. High compressive residual stresses occurred in all layers. Stresses of −1.4 GPa were measured also in coatings deposited without bias. With increasing bias the residual stress increased to −4.5 GPa at 150 V. The amount of stress generated was dependent on the level of argon content but not on the amount of boron (Fig. 6). Therefore, Ar incorporation can be regarded as the main reason for stress formation. Unfortunately, the adhesion of coatings deposited with higher bias voltages on silicon substrates was not sufficient. Partial delaminations took place and, therefore, the measurement of residual stresses by bending was impossible. Strong texture, high compressive stresses and low intensities cause great problems for the estimation of the correct lattice constants. Nevertheless, measured values can be compared with one another. The following lattice parameters of TiB phase were calculated: 2 0 V bias: a=0.3051±0.0011 nm, c=0.3193±0.0012 nm; −200 V bias: a=0.3044±0.0010 nm, c=0.3224± 0.0011 nm.
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The errors were relatively great as a consequence of difficulties mentioned above. The measured values were not stress corrected and, therefore, they are not comparable with bulk material directly (bulk: a=0.303034 nm, c=0.322953 nm [6 ]). However, the lattice parameter a decreased and the parameter c increased with rising bias corresponding to higher residual stresses, meaning that the hexagonal lattice was stretched in (001) direction. It is not clear where argon and the excess boron are placed. In the hexagonal TiB lattice, incorporation of 2 interstitial atoms is only possible in tetrahedral holes. However, the misfit of hole and argon radii is about 60% and Ar incorporation is unsuccessful. The incorporation of boron in tetrahedral holes can also be excluded. Substitution of Ti atoms by argon or boron atoms is also improbable owing to the large difference in the radii (r =0.14 nm, r =0.10 nm, r =0.09 nm) and Ti Ar B would produce tensile stresses. In the TiN lattice Ar atoms substitute N atoms under generation of compressive stresses [7]. The same result should be possible in TiB when Ar atoms occupy boron 2 sites, but this reflection does not agree with the measured high boron content.
3.4. Mechanical properties Direct measurement of conventional Vickers hardness is impossible because the layers are very brittle and many cracks are formed in and around all indentations. Therefore, the diagonals of indentations cannot be determined correctly. For that reason, hardness testing under load is used where the indentation depth of a Vickers indenter is measured. Commonly, in thin layers the direct determination of layer hardness is impossible because the indentation depth is larger than the layer thickness. It is possible to calculate layer hardness without the influence of substrate hardness, using a method described in Ref. [8]. In this case, the measured compound hardness values estimated at different loads are expressed as a function of the ratio of the layer thickness to the indentation depth. The layer hardness can be calculated from the slope of the corresponding curve. Using this method the following layer hardness values were determined: 0 V bias: HU=31 500 N mm−2; −200 V bias: HU= 34 300 N mm−2. The results show that the high compressive residual stresses increase the hardness of TiB layers. Similar 2 results on the influence of residual stresses have been presented previously [9,10]. However, the measured hardness values of TiB fluctuate between 3000 and 2 6700 HV. This is caused by the difficulties of thin-layer hardness measurements which are described above and more in detail in Ref. [8].
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Fig. 4. X-Ray diffraction patterns.
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Fig. 5. Quantitative pole figures.
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Table 1 Integral intensities (I ) and integral breadth of TiB layers deposited 2 with various bias voltages Bias voltage (V )
I (deg s−1)
B (°2H)
0 −50 −200
1106 286 8
0.32 0.39 0.91
stresses generated as a sequence of soft substrate material displacements during penetration of the stylus. Higher compressive stresses prevent cohesive cracking under the stylus and favour spalling or buckling in front of the indenter. Fig. 7 gives typical examples of different failure modes. Our results are in contrast to previous measurements [3], where the critical load decreased with increasing bias. However, for further discussion more information on failure modes and coating properties is required.
3.5. Annealing
Fig. 6. Residual stresses and argon content in dependence on bias voltage.
The compressive residual stresses also improve the adhesion of layers and influence the failure modes during scratch testing. The critical load of adhesive failure increases with rising bias from 4 N at 0 V to 19 N at −200 V. In general, the measured critical load values are very low, because of the soft substrate. The substrate hardness amounts only 380 HV 0.1. Other authors [4] have measured 50–90 N for TiB layers on cemented 2 carbide. The failure of coatings deposited with lower compressive stresses (0 V bias) takes place with long cracks starting in the middle and small shell-shaped delaminations at the edges of scratch tracks. Samples with higher compressive stresses exhibit only small cracks at the track edges using the same normal force. These layers fail by upwelling and spalling in front of the stylus at higher forces. The first type of cracks and small delaminations are caused by tensile
After annealing at 400 °C the chemical composition and the microstructure (phases and texture) are not modified and the critical load of adhesive failure is independent of the bias voltage. The values of all investigated samples amount to only 3–4 N and, thus, they are placed in the range of 0 V bias samples before annealing. It is supposed that in this state the compressive residual stresses produced during deposition should be relaxed. Therefore, these stresses do not improve the adhesion. Nevertheless, the failure mode is the same as before annealing, which may indicate that the failure mode is also influenced by morphology (smaller grain size with increasing bias). During heating to 400 °C tensile stresses up to 1.3 GPa are produced, which are caused by the difference between the thermal expansion coefficients of iron and TiB 2 (a =12 · 10−6 K−1, a =7.8 10−6 K−1), but cracks Fe TiB2 could not be observed in the layers. Accordingly, the layers are able to compensate these stresses elastically. After 800 °C annealing the morphology of all coatings is modified. The surface is divided by many cracks because tensile stresses grow to 2.6 GPa during heating. The crack pattern is generated periodically. Furthermore, an upwelling of the surface between the cracks can be observed. The cross-sections show three layers: two (I, II ) in the former coating and one (III ) with large and lengthy grains and large pores under the former interface between layer and substrate (Fig. 8). The thickness of layer I and II is in the same range as the TiB layer before annealing. Layer III is up to 2 6 mm thick. The GDOES quantitative concentration–depth profiles exhibit a new distribution of B, C and Fe, i.e. distant diffusion occurs (Fig. 9). Boron diffuses up to 10 mm into the steel substrate. The profiles can also be divided in three parts: layer I, on the surface with high boron content; layer II, below, with a maximum of boron and titanium and a minimum of boron content; and layer III, in the former steel substrate with a second maximum of boron content. The X-ray phase diagram of symmetric X-ray diffraction also contains TiB and three new phases, TiC, FeTi 2
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Fig. 7. Failure modes during scratch testing (above, low compressive stresses; below, high compressive stresses).
and Fe B. The X-ray phase analyses with glancing 2 incidence at different angles prove that the phase Fe B 2 only occurs under the former interface. From the concentration–depth profiles and the X-ray phase analyses it can be concluded that the first layer with the highest boron content should consist mainly of TiB . It can be assumed that TiC occurs in the second 2 layer where a minimum of boron and a maximum of carbon content exist. Beside this, the formation of FeTi is possible. In the
upper range of steel substrate the ferrite is completely transformed in Fe B, which is connected with new grain 2 formation and growth. The phase Fe C is modified in 3 Fe B and free carbon which diffuses in the former 2 Ti B layer and combines with free titanium. Therefore, 2 from surface to substrate the following phases could be formed: layer I, TiB ; layer II, TiC, FeTi, also possibly 2 TiB ; layer III, Fe B. 2 2 The generation of the new phases is a consequence of the free enthalpy reduction. The following chemical
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Fig. 8. SEM imaging of fracture cross-section after annealing (above, SE contrast; below, BSE contrast).
Fig. 9. GDOES concentration–depth profile after annealing (I–III, layer thickness measured at fractured samples, see Fig. 8).
reactions take place: 3TiB +6Fe+Fe CTiB +TiC+TiFe+4Fe B 2 3 2 2 S°G=−302 kJ K mol−1 S°G=−384 kJ K mol−1 D°G=−82 kJ K mol−1. The layers are also very brittle and not suitable for application.
4. Conclusions In this investigation the structure of TiB layers 2 deposited with magnetron sputtering at −200 V bias
was nanocrystalline (grain diameters 5–40 nm) and not amorphous. The boron content was higher than in the stoichiometric one, but no crystalline boron phases could be detected by XRD or TEM. The occurrence of the Ti B phase 3 4 could not explain the high boron content. It should be excluded that boron atoms are incorporated in interstitial holes or on regular Ti atom sites. The authors of Ref. [4] supposed that the excess B atoms are incorporated in the grain boundaries of nanocrystals. The incorporation of Ar atoms in the hexagonal TiB lattice is regarded as the main reason for the 2 generation of high compressive stresses in sputtered TiB coatings. It is unclear on which sites in the lattice 2 the argon atoms are placed. Higher compressive stresses increase the adhesion and the hardness of TiB layers. The value of critical load 2 should be improved by harder substrates. TiB layers react with steel substrates at T>400 °C. 2 A diffusion barrier layer should be used for applications at higher temperatures, but crack formation caused by various thermal coefficients will also occur. These features limit the application temperature of these layers.
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by means of X-ray diffractometry and Auger electron spectroscopy, Mat. Sci. Engng A139 (1991) 259–263. ¨ C. Mitterer, M. Rauter, P. Rodhammer, Sputter deposition of ultrahard coatings within the system Ti–B–C–N, Surf. Coat. Technol. 41 (1990) 351–363. T. Shikama, et al., Deposition of TiB films by a co-sputtering 2 method, Thin Solid Films 156 (1988) 287–293. Powder Diffraction File PDF-2 Database Sets 1–42, International Centre for Diffraction Data (JCPDS), 1992. R. Wiedemann, H. Oettel, Macroscopic and microscopic stresses in nitride hard coatings, in: V. Hauk et al., Residual Stresses; DGM Informationsges, 1993, 673–681.
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¨ [8] A. Kirsten, R. Wiedemann, H. Oettel, Ermittlung der Schichtharte ¨ ¨ aus gemesenen Verbundhartewerten, Harterei-Techn. Mitt. 48 (1993) 324–329. [9] W. Herr, B. Matthes, E. Broszeit, K.H. Kloos, Fundamental properties and wear resistance of r.f. sputtered TiB and Ti(B,N ) 2 coatings, Mat. Sci. Engng A140 (1991) 616–624. [10] O. Knotek, F. Jungblut, R. Breitenbach, Magnetron-sputtered superhard coatings within the system Ti–B–C–N, Vacuum 41 (1990) 2184–2186.