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Fatigue behavior of (graded) (Ti, Al)N-coated 1Cr11Ni2W2MoV stainless steel at high temperature
Surface & Coatings Technology 204 (2010) 2417–2423
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Fatigue behavior of (graded) (Ti, Al)N-coated 1Cr11Ni2W2MoV stainless steel at high temperature Li Xin a,⁎, Ping Liu a, Changjie Feng b, Shenglong Zhu a, Fuhui Wang a a b
State key Laboratory for Corrosion and Protection of Metals, Institute of Metal Research, Chinese Academy of Sciences, Wencui Road 62, Shenyang, 110016, China School of Materials Science and Engineering, Nanchang Hangkong University, Nanchang, 330063, China
a r t i c l e
i n f o
Article history: Received 30 April 2009 Accepted in revised form 8 January 2010 Available online 25 January 2010 Keywords: (Graded ) (Ti, Al)N coating Arc ion plating High-cycle fatigue High temperature Stainless steel
a b s t r a c t The high-cycle fatigue properties of graded (Ti, Al)N- and Ti0.7Al0.3N-coated 1Cr11Ni2W2MoV at 500 °C have been investigated using a rotating bending fatigue testing machine. The results show that fatigue strength and life of 1Cr11Ni2W2MoV stainless steel were apparently increased by the presence of Ti0.7Al0.3N coating. The fatigue life and strength was also improved to some extent by the presence of the graded coating at higher stress levels (475 MPa–525 MPa). The fracture morphologies were analyzed by SEM. It is concluded that the presence of the coatings, which were well adhered and have high compressive stress, restricted plastic deformation of the substrate, thus improved the fatigue properties of the steel. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Hard coatings characterized by high hardness and wear resistance as well as good corrosion resistance are more and more widely used in the aircraft, transports, oil and biomedical industry in recent years [1]. In addition to the wear and the corrosion properties of the coatings, the main focuses of the application, fatigue behaviors of the coating– substrate systems have attracted more attention and have been widely studied in the last few years [1–17] due to applications of the coatings on connecting rods, crankshafts, gear wheel and turbine rotors where the fatigue property is an important parameter. The properties of the coating, such as composition, defects, surface roughness, stress level and mechanical properties, will affect the fatigue behavior of the substrate. The thin hard coatings deposited by PVD usually have positive effect on the fatigue properties of the substrate at room temperature. The fatigue properties of TiN-, TiCN- and TiAlN-coated Cr–Mo–V steel [2], TiN- and ZrN-coated 316L steel [3–5], TiN-coated AISI 1045 steel [6], (graded) (Ti, Al)N-coated 1Cr11Ni2W2MoV [7] and CrN-coated 15NiCr13 [1] have been investigated and the results show that the presence of the coatings improves the fatigue strength of the steel. The compressive residual stress, the high mechanical strength of the film and its excellent adhesion to the substrate are considered to be related to the improved performance. The fatigue properties of TiNcoated titanium have also been investigated and similar results obtained [1]. One exception is the ZrN-coated AA7075-T6 aluminum
⁎ Corresponding author. Tel.: +86 24 23887796; fax: +86 24 23893624. E-mail address: [email protected] (L. Xin). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.01.013
alloy, whose fatigue properties were inferior to that of the substrate aluminum alloy [8]. Numerical models predicting the fatigue life procedure of the PVD coated components have been introduced [1,9]. The fatigue behavior of duplex treated steel coated with PVD Cr(C, N) [10] and PACVD TiC xN1−x and Ti1−xAlxN coating [11] have also been studied. It is found that whether the coatings will have significant effect on the fatigue strength of the pre-treated substrate depend on the compressive residual stresses of the films. The type and stoichiometry of the coating will influence the fatigue behavior. On the other hand, the hard coating with tensile stress such as eletrodeposited chromium plating [12–14] usually decreases the fatigue properties of the substrate. The effects of subsequent treatment like shot peening used for introduction of residual stress on the fatigue strength of the chromium plating steel [12] and aluminum alloy [14] were studied. In addition, WC–Co–(Cr), cleaner alternatives to the hard chromium plating, were applied by HVOF on steel and aluminum alloy and fatigue behaviors of the coating– substrate systems were extensively investigated [12,13,15,16]. If a coating–substrate system is to be used for aero-engine compressor blades, knowledge of its fatigue properties at elevated temperature is also essential. However, there have been relatively few studies [17] about the effect of hard coatings on fatigue characteristics at elevated temperature, where the stress level in the coating, and the mechanical properties of the coating and the substrate, will be different from those at room temperature. Furthermore the defects and composition of the coating may change due to high-temperature oxidation. Thus the fatigue properties of the coating–substrate system may not be the same as those at room temperature. (Ti, Al)N coatings, which exhibit good wear resistance, extreme micro-hardness and excellent oxidation and corrosion resistance,
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have proved to be good candidates for the corrosion protection of the stainless steel for aero-engine compressor blades [7,18–20]. In the previous work [7,18–20], mechanical properties including microhardness, wear and adhesion between coating and substrate, depth distribution of residual stress, and oxidation and corrosion resistance of the (graded) (Ti, Al)N coating and fatigue properties of (graded) (Ti, Al)N-coated stainless steel at room temperature were studied. In the present study, the fatigue properties of graded (Ti, Al)N- and Ti0.7Al0.3N-coated 1Cr11Ni2W2MoV at 500 °C have been investigated. 2. Experimental procedure 2.1. Sample preparation and coating process The substrate material was 1Cr–11Ni–2W–2Mo–V stainless steel. Test specimens were cut from forged rods, which had been homogenized at 1010 °C for 1 h, oil quenched, and then tempered at 570 °C for 1.5 h. The specimens for the fatigue tests were machined as shown in Fig. 1. The surfaces of the specimens were polished and the average roughness at the constant cross-sectional gauge length of the specimens was approximately Ra = 0.35 μm. Ti0.7Al0.3N and graded (Ti, Al)N were deposited on the specimens by arc ion plating in a coating unit (DH-4, China). The details of the deposition process have been reported elsewhere [18–20]. The temperature of the deposition chamber was controlled at 230 °C. The deposition time was 60–70 min. In the case of the Ti0.7Al0.3Ncoated specimens TiN was first deposited to act as a buffer layer. The thickness of the Ti0.7Al0.3N and graded (Ti, Al)N coating was 2.7 μm and 2.9 μm respectively. The average roughness of the coated specimens was almost as same as that of the uncoated specimens. 2.2. Fatigue test The fatigue tests were carried out using a rotating bending fatigue testing machine PWC510WC with a rotation speed of 5000 rpm and
R = −1 at 500 °C in room air. The three groups of specimens prepared to determine the fatigue limit (the infinite life was specified at a number of 1 × 107 cycles according Chinese standard GB 2107-80) and obtain the S–N curves in the fatigue test were listed below: 13 smooth specimens of bare steel; 15 smooth specimens with Ti0.7Al0.3N coating; 15 smooth specimens with graded (Ti, Al)N coating. The fatigue limit was determined by means of the staircase method employing a step of 25 MPa and 8 specimens for each condition, and calculated according the following equation: σ=
1 n ∑ σ: n i=1 i
ð1Þ
Here n is the number of the efficient experiment and σi is the alternating stress. After the fatigue limit was determined, the fatigue test was performed at higher maximum alternating stresses than the fatigue limit and the increment was 25 MPa each time. Two or three specimens were tested at each alternating stress. The fracture appearance, surface and cross-section of the specimens were observed by SEM. 3. Results 3.1. Fatigue strength and life Tables 1–3 present the data concerning the determination of the fatigue limit of the uncoated and coated specimens. It can be seen that at 500 °C in air, the fatigue limit was 450 MPa for uncoated and graded (Ti, Al)N-coated specimens, and 494 MPa for Ti0.7Al0.3Ncoated specimens. Thus the deposition of Ti0.7Al0.3N coating gave an increase of around 10% in fatigue limit. The deposition of the graded (Ti, Al)N coating did not give any increase in fatigue limit.
Fig. 1. Geometry of the rotary bending fatigue test specimen; dimension in mm.
L. Xin et al. / Surface & Coatings Technology 204 (2010) 2417–2423
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Table 1 Experimental results for determining the fatigue limit of the uncoated samples. Sample
Alternating stress (MPa)
Number of cycles
1 2 3 4 5 6 7 8 Fatigue limit
450 475 450 475 450 425 450 425 450
11,700,000 7,335,000 11,500,000 255,000 1,740,000 10,000,000 9,500,000 10,000,000
Table 2 Experimental results for determining the fatigue limit of the graded (Ti, Al)N-coated samples. Sample
Alternating stress (MPa)
Number of cycles
1 2 3 4 5 6 7 8 Fatigue limit
450 425 450 475 450 475 450 425 450
4,500,000 10,000,000 10,000,000 9,000,000 10,000,000 4,550,000 8,170,000 10,000,000
Fig. 2. The S–N curves of the substrate, graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens at 500 °C in air. The data used to determine the fatigue limit are indicated by the solid symbols.
Table 4 Parameters involved in the Basquin relationship for the conditions tested.
Table 3 Experimental results for determining the fatigue limit of the Ti0.7Al0.3N-coated samples. Sample
Alternating stress (MPa)
Number of cycles
1 2 3 4 5 6 7 8 Fatigue limit
475 500 525 500 475 500 475 500 494
10,000,000 10,000,000 600,000 9,385,000 10,000,000 3,285,000 10,000,000 575,000
−m
;
A (MPa)
m
741.1 951.4 850.7
0.031 0.046 0.034
3.2. Fractographic analysis
Fig. 2 presents all the data obtained in the fatigue test. Basquin equation, S = ANf
Condition Substrate Graded (Ti, Al)N Ti0.7Al0.3N
After the fatigue test at 500 °C, the steel was slightly oxidized, both on the surface of the specimen and the fracture surface. Fig. 4a illustrates the fracture surface of the uncoated specimen tested at 450 MPa. Fig. 4b is the magnified view of Fig. 4a at the crack initiation site. The photomicrographs show that fracture of the sample occurred due to propagation of a crack which originated from the surface of the sample. A similar conclusion can be drawn from observation of the fracture surfaces of specimens tested at higher stress levels.
ð2Þ
was employed to describe the relationships of the number of cycles to fracture and the alternating stresses of the uncoated and coated specimens, and the curves with internals indicating 95% confidence were plotted in Fig. 2. It can be seen that the fatigue strength was also improved to some extent by the graded coating at higher stress levels (475 MPa–525 MPa). The values of the fatigue strength coefficient A and the fatigue strength exponent m in Basquin relationship were summarized in Table 4. The increase in fatigue life can be computed as: %increase =
Nfcoating −Nfsubstrate Nfsubstrate
× 100%
ð3Þ
where the number of cycles to fracture for each condition is calculated from Eq. (3), employing the values given in Table 4. The results were presented in Fig. 3. For the Ti0.7Al0.3N coating, the increase in fatigue life ranges between 1785 and 2706%, when testing is conducted at stresses between 500 and 575 MPa. For the graded (Ti, Al)N coating, the increase in fatigue life ranges between 112 and 507%, when testing is conducted at stresses between 475 and 525 MPa.
Fig. 3. Change in the percentage of increase in fatigue life for Ti0.7Al0.3N- and graded (Ti, Al)N-coated specimen.
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Fig. 4. (a) SEM fatigue fracture surface and (b) magnified view of the crack initiation site in a for uncoated steel tested at 450 MPa.
Fig. 5. (a) SEM fatigue fracture surface and (b) magnified view of the crack initiation site in a for graded (Ti, Al)N-coated steel tested at 450 MPa.
Although the beneficial effect of the presence of the graded (Ti, Al)N on the fatigue life and strength of the steel is not as significant as that of Ti0.7Al0.3N coating, the fracture surfaces of both kinds of specimens show similar characteristics. The fracture surface appearances of the graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens tested at 450 MPa and 500 MPa are shown in Figs. 5 and 6. It can be seen that for the coated specimen, the fatigue crack propagated from the surface of the substrate into the interior, and eventually caused the failure of the specimen, as shown in
Figs. 5a–b and 6a–b. For the coated specimens tested at higher stress levels, similar characteristics were observed. The coatings remain well adherent to the substrate during the fatigue tests at lower stress level (450 MPa–525 MPa). At higher stress (575 MPa), extensive fracture and spallation of Ti0.7Al0.3N coating occurred. In order to obtain more information about the fatigue crack initiation, surfaces of the graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens, which were tested at 450 MPa and 500 MPa respectively and did not fail over a
Fig. 6. (a) SEM fatigue fracture surface and (b) magnified view of the crack initiation site in a for Ti0.7Al0.3N-coated steel tested 500 MPa.
L. Xin et al. / Surface & Coatings Technology 204 (2010) 2417–2423
Fig. 7. SEM surface morphology of a spalled area on a Ti0.7Al0.3N-coated specimen tested at 500 MPa for 1 × 107 cycles.
1×107 cycle test, were observed. For both kinds of specimens, spallation of the coatings, immediately adjacent to the machining mark, occurred occasionally. The typical view of the spalled areas is shown in Fig. 7. It can be seen that the spallation occurred not along the interface, instead with
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an inclination to it, and the substrate was exposed only at the edge of the groove (machining mark). In some cases, the crack grew within the coatings and led to the spallation, and the substrate was not exposed. The macro-particles which are inevitably deposited during the coating process are also visible, and seem to have negligible effect on the coating spallation and fatigue crack initiation. A fatigue crack, initiated from the substrate surface where the coating had spalled (the edge of the groove), was found in a graded (Ti, Al)N-coated specimen, as shown in Fig. 8a–b. The coating was stripped away by a chemical stripping solution, which did not react with the steel substrate [7], so that more information about the crack can be obtained. As shown in Fig. 8c, the length of the crack is about 25 μm longer than that observed before the coating stripped away, indicating crack propagation in the substrate underneath the coating (Fig. 8d). The cross-sections of specimens were also prepared normal to the surface, taken along the axis of the specimen. Fig. 9a shows cracks within the coating and Fig. 9b shows the cross-section of a typical spalled area shown in Fig. 7. It indicates that the cracks may initiate within the coating or at the coating/substrate interface, and propagate within the coating rather than along the interface, eventually reach the coating surface and lead to the spallation. It also indicates that the bond strength between the coating and the substrate is quite strong. It is worth to mention that, for the graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens tested at 450 MPa and 500 MPa for 1×107 cycles respectively and did not fail, the spallation of the coating, as shown in Fig. 7, was observed in all the
Fig. 8. (a) SEM surface morphology of a crack originated from the substrate surface where the coating had spalled on a graded (Ti, Al)N-coated specimen tested at 450 MPa for 1 × 107 cycles; (b) magnified view of a; (c) corresponding BSE view of a after coating stripped; and (d) magnified view of the square area in c.
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Fig. 9. BSE cross-sections normal to the specimen surface, taken along the axis of coated specimens. (a) graded (Ti, Al)N-coated specimen tested at 450 MPa for 1 × 107 cycles and (b) Ti0.7Al0.3N-coated specimen tested at 500 MPa for 1 × 107 cycles.
specimens, but the occurrence of fatigue crack, as shown in Fig. 8, was only observed in one specimen. Cross-sections of the failed specimens in the fatigue test were also prepared. Fig. 10 is the cross-section of a Ti0.7Al0.3N-coated specimen tested at 500 MPa and failed at 3.29 × 106 cycles. It can be seen the fatigue crack originated from the substrate surface had propagated into the substrate even when through-scale crack has not formed in the coating well adherent to the substrate. 4. Discussion From Fig. 2, it can be seen the standard deviation of the number of cycles to fracture at various maximum alternating stresses is quite large for uncoated and coated specimens. This may be related with the surface roughness of the specimen. It has been measured that the average roughness at the constant cross-sectional gauge length of the specimens was approximately Ra = 0.35 μm after polishing, but in some areas Ra could reach to 0.67 μm (Fig. 9). After the deposition of the coating, the average surface roughness measured was almost as same as that of the uncoated specimen, and from the cross-sections of the coated specimens (Fig. 9), it can also be observed that the deposition of the coatings did not change the surface roughness obviously. However, most of the experimental data were in or near to the intervals of 95% confidence after S–N curves were plotted employing Basquin relationship. So, from the statistical point of view, it can be included the fatigue life and strength was improved apparently by Ti0.7Al0.3N coating at all the
Fig. 10. BSE cross-section normal to the specimen surface, taken along the axis of Ti0.7Al0.3N-coated specimen tested at 500 MPa and failed at 3.29 × 106 cycles.
testing stress levels, and improved to some extent by the graded coating at higher stress levels (475 MPa–525 MPa). The deposition of Ti0.7Al0.3N coating also gave an increase of around 10% in fatigue limit. It has been measured that the average residual compressive stresses in graded (Ti, Al)N and Ti0.7Al0.3N coatings were around 4 GPa at room temperature [7]. At 500 °C, the compressive stresses in the coatings will be lower than that at room temperature due to the thermal expansion discrepancy between the coating and the substrate, and the stress relaxation (SR) in the coating resulting from the creep deformation of the substrate near the interface at high temperature [17]. Thermal stress can be quantified by the following equation [21]: σth =
Ef ðα −αs ÞðTt −Tr Þ: 1−νf f
ð4Þ
In the above equation, Ef and νf are Young's modulus and Poisson ratio of film, respectively; αf and αs are thermal expansion coefficients of film and substrate; Tt is the temperature of the fatigue test and Tr room temperature. If we use αs = 11.7 × 10− 6 °C− 1, αTiN = 9.4 × 10− 6 °C− 1, αTiAlN = 7.5 × 10− 6 °C− 1, ν = 0.2, and E = 450 GPa [22] in Eq.(4), σth in TiN is about 0.61 GPa and in Ti0.5Al0.5N 1.12 GPa. Suh et al. [17] measured the stress relaxation of TiAlN coating at high temperature and found that the decrement of the residual stress in the coating was around 0.5 GPa after SR treatment at 500 °C. Taking the above two factors into account, the average stresses in the coatings were still compressive at 500 °C. The micro-hardness was 25.1 GPa for the graded coating [19] and 23.5 GPa for Ti0.7Al0.3N coating [20] respectively at room temperature. The difference of the two coatings in micro-hardness was not significant. Cracking and spallation of the coatings occasionally observed on the coated specimens which did not fail in the fatigue test is due to the propagation of the preexisting defects within the coatings or at the coating/substrate interface under the cyclic loading at the sites of stress concentration due to localized change of curvature, for example, at the machining marks. The spallation of the coatings occurred not along the interface instead with an inclination indicated that the coating–substrate interface is quite strong and more resistant to the rupture than is the coating itself. The reason for the cracks in the coatings having not penetrated into the substrate may lie in that the propagation directions of the cracks were not perpendicular to the interface. The fatigue crack was observed to initiate at the substrate surface where the coating had spalled under the cyclic loading on the coated specimen which did not fail in the fatigue test. Furthermore, the fatigue cracks had initiated and propagated into the substrate even
L. Xin et al. / Surface & Coatings Technology 204 (2010) 2417–2423
when the through-scale crack has not formed in the coating which was well adherent to the substrate on the fractured specimens. At 500 °C, the elastic modulus (142 GPa) of 1Cr11Ni2W2MoV steel decreased significantly compared with that at room temperature (196 GPa). The substrate became softer and easier to deform under the alternating stress. The presence of the coatings with high compressive stress will restrict the deformation of the substrate. At the place where the coating had spalled, the restriction to the substrate diminished and the fatigue crack was easier to initiate. At the place where the coating maintained adherent to the substrate, when the plastic deformation of the substrate was too large to be accommodated by the coating, fracture of the coating occurred, and fatigue crack initiated and propagated into the coating and the substrate. The fatigue cracks even had propagated into the substrate when the through-scale crack has not formed in the coating. The compressive coating may retard the growth of surface crack tip to some extent (Fig. 8). Baragetti [1] found that the specified crack propagation depth of CrN-coated steel was deeper than that of the substrate steel and it was the high residual stresses at the surface that retarded the opening of the crack. Tests to clarify the effect of the coatings on crack growth rate in the substrate have been carried out by Kim et al. [2] and the results showed that the presence of the coating layer did not significantly retard crack growth. Therefore it is reasonable to think the presence of the coating will retard the growth of the crack at the very beginning and afterwards its contribution to the restriction of crack growth will be limited. So it can be concluded that main reason the presence of the graded (Ti, Al)N and Ti0.7Al0.3N coatings improving the fatigue life and strength of the steel is that the coatings, which are well adhered and have high residual stress, restricted the deformation of the substrate and delayed the fatigue crack initiation. During fatigue tests at 500 °C, the uncoated specimens were slightly oxidized, which promoted the initiation of the fatigue cracks, but the oxidation of the coatings was very slight, and seems to have negligible effect on the fatigue initiation. This may also account for the improvement of the fatigue properties by the presence of the coatings. The beneficial effect of the presence of the graded (Ti, Al)N on the fatigue life and the fatigue strength of the steel is not as significant as that of Ti0.7Al0.3N coating, which may be related to the difference on the compressive stress in the coatings [7]. For the graded coating, Al content in the coating increased gradually from substrate/coating interface to coating surface, and Ti concentration changed in the opposite direction, thus the inner layer rich in TiN and outer layer rich in Ti0.5Al0.5N. Compared with Ti0.7Al0.3N coating, the stress maximum value in the graded coating was higher and nearer to the coating surface, so the graded coating may be easier to crack and spall under the alternating stress and the restraint to the substrate decreased [7].
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5. Conclusions 1. At 500 °C, the fatigue life and strength of 1Cr11Ni2W2MoV stainless steel were apparently increased by the presence of Ti0.7Al0.3N coating. The fatigue life and strength of the steel was also improved to some extent by the presence of the graded coating at higher stress levels (475 MPa–525 MPa). 2. The presence of the coatings, which were well adhered and have high compressive stress, restricted the plastic deformation of the substrate, thus improved the fatigue properties of the steel.
Acknowledgement This work was supported by the National Nature Foundation of China under grant no. 50671110. The authors greatly appreciate the very helpful discussion with Professors Guangping Zhang, Weicheng Yu, Suhua Ai and Huihe Su about the fatigue results, and wish to thank Mr. Shichen Wang for the coating deposition.
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Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Fatigue behavior of (graded) (Ti, Al)N-coated 1Cr11Ni2W2MoV stainless steel at high temperature Li Xin a,⁎, Ping Liu a, Changjie Feng b, Shenglong Zhu a, Fuhui Wang a a b
State key Laboratory for Corrosion and Protection of Metals, Institute of Metal Research, Chinese Academy of Sciences, Wencui Road 62, Shenyang, 110016, China School of Materials Science and Engineering, Nanchang Hangkong University, Nanchang, 330063, China
a r t i c l e
i n f o
Article history: Received 30 April 2009 Accepted in revised form 8 January 2010 Available online 25 January 2010 Keywords: (Graded ) (Ti, Al)N coating Arc ion plating High-cycle fatigue High temperature Stainless steel
a b s t r a c t The high-cycle fatigue properties of graded (Ti, Al)N- and Ti0.7Al0.3N-coated 1Cr11Ni2W2MoV at 500 °C have been investigated using a rotating bending fatigue testing machine. The results show that fatigue strength and life of 1Cr11Ni2W2MoV stainless steel were apparently increased by the presence of Ti0.7Al0.3N coating. The fatigue life and strength was also improved to some extent by the presence of the graded coating at higher stress levels (475 MPa–525 MPa). The fracture morphologies were analyzed by SEM. It is concluded that the presence of the coatings, which were well adhered and have high compressive stress, restricted plastic deformation of the substrate, thus improved the fatigue properties of the steel. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Hard coatings characterized by high hardness and wear resistance as well as good corrosion resistance are more and more widely used in the aircraft, transports, oil and biomedical industry in recent years [1]. In addition to the wear and the corrosion properties of the coatings, the main focuses of the application, fatigue behaviors of the coating– substrate systems have attracted more attention and have been widely studied in the last few years [1–17] due to applications of the coatings on connecting rods, crankshafts, gear wheel and turbine rotors where the fatigue property is an important parameter. The properties of the coating, such as composition, defects, surface roughness, stress level and mechanical properties, will affect the fatigue behavior of the substrate. The thin hard coatings deposited by PVD usually have positive effect on the fatigue properties of the substrate at room temperature. The fatigue properties of TiN-, TiCN- and TiAlN-coated Cr–Mo–V steel [2], TiN- and ZrN-coated 316L steel [3–5], TiN-coated AISI 1045 steel [6], (graded) (Ti, Al)N-coated 1Cr11Ni2W2MoV [7] and CrN-coated 15NiCr13 [1] have been investigated and the results show that the presence of the coatings improves the fatigue strength of the steel. The compressive residual stress, the high mechanical strength of the film and its excellent adhesion to the substrate are considered to be related to the improved performance. The fatigue properties of TiNcoated titanium have also been investigated and similar results obtained [1]. One exception is the ZrN-coated AA7075-T6 aluminum
⁎ Corresponding author. Tel.: +86 24 23887796; fax: +86 24 23893624. E-mail address: [email protected] (L. Xin). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.01.013
alloy, whose fatigue properties were inferior to that of the substrate aluminum alloy [8]. Numerical models predicting the fatigue life procedure of the PVD coated components have been introduced [1,9]. The fatigue behavior of duplex treated steel coated with PVD Cr(C, N) [10] and PACVD TiC xN1−x and Ti1−xAlxN coating [11] have also been studied. It is found that whether the coatings will have significant effect on the fatigue strength of the pre-treated substrate depend on the compressive residual stresses of the films. The type and stoichiometry of the coating will influence the fatigue behavior. On the other hand, the hard coating with tensile stress such as eletrodeposited chromium plating [12–14] usually decreases the fatigue properties of the substrate. The effects of subsequent treatment like shot peening used for introduction of residual stress on the fatigue strength of the chromium plating steel [12] and aluminum alloy [14] were studied. In addition, WC–Co–(Cr), cleaner alternatives to the hard chromium plating, were applied by HVOF on steel and aluminum alloy and fatigue behaviors of the coating– substrate systems were extensively investigated [12,13,15,16]. If a coating–substrate system is to be used for aero-engine compressor blades, knowledge of its fatigue properties at elevated temperature is also essential. However, there have been relatively few studies [17] about the effect of hard coatings on fatigue characteristics at elevated temperature, where the stress level in the coating, and the mechanical properties of the coating and the substrate, will be different from those at room temperature. Furthermore the defects and composition of the coating may change due to high-temperature oxidation. Thus the fatigue properties of the coating–substrate system may not be the same as those at room temperature. (Ti, Al)N coatings, which exhibit good wear resistance, extreme micro-hardness and excellent oxidation and corrosion resistance,
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have proved to be good candidates for the corrosion protection of the stainless steel for aero-engine compressor blades [7,18–20]. In the previous work [7,18–20], mechanical properties including microhardness, wear and adhesion between coating and substrate, depth distribution of residual stress, and oxidation and corrosion resistance of the (graded) (Ti, Al)N coating and fatigue properties of (graded) (Ti, Al)N-coated stainless steel at room temperature were studied. In the present study, the fatigue properties of graded (Ti, Al)N- and Ti0.7Al0.3N-coated 1Cr11Ni2W2MoV at 500 °C have been investigated. 2. Experimental procedure 2.1. Sample preparation and coating process The substrate material was 1Cr–11Ni–2W–2Mo–V stainless steel. Test specimens were cut from forged rods, which had been homogenized at 1010 °C for 1 h, oil quenched, and then tempered at 570 °C for 1.5 h. The specimens for the fatigue tests were machined as shown in Fig. 1. The surfaces of the specimens were polished and the average roughness at the constant cross-sectional gauge length of the specimens was approximately Ra = 0.35 μm. Ti0.7Al0.3N and graded (Ti, Al)N were deposited on the specimens by arc ion plating in a coating unit (DH-4, China). The details of the deposition process have been reported elsewhere [18–20]. The temperature of the deposition chamber was controlled at 230 °C. The deposition time was 60–70 min. In the case of the Ti0.7Al0.3Ncoated specimens TiN was first deposited to act as a buffer layer. The thickness of the Ti0.7Al0.3N and graded (Ti, Al)N coating was 2.7 μm and 2.9 μm respectively. The average roughness of the coated specimens was almost as same as that of the uncoated specimens. 2.2. Fatigue test The fatigue tests were carried out using a rotating bending fatigue testing machine PWC510WC with a rotation speed of 5000 rpm and
R = −1 at 500 °C in room air. The three groups of specimens prepared to determine the fatigue limit (the infinite life was specified at a number of 1 × 107 cycles according Chinese standard GB 2107-80) and obtain the S–N curves in the fatigue test were listed below: 13 smooth specimens of bare steel; 15 smooth specimens with Ti0.7Al0.3N coating; 15 smooth specimens with graded (Ti, Al)N coating. The fatigue limit was determined by means of the staircase method employing a step of 25 MPa and 8 specimens for each condition, and calculated according the following equation: σ=
1 n ∑ σ: n i=1 i
ð1Þ
Here n is the number of the efficient experiment and σi is the alternating stress. After the fatigue limit was determined, the fatigue test was performed at higher maximum alternating stresses than the fatigue limit and the increment was 25 MPa each time. Two or three specimens were tested at each alternating stress. The fracture appearance, surface and cross-section of the specimens were observed by SEM. 3. Results 3.1. Fatigue strength and life Tables 1–3 present the data concerning the determination of the fatigue limit of the uncoated and coated specimens. It can be seen that at 500 °C in air, the fatigue limit was 450 MPa for uncoated and graded (Ti, Al)N-coated specimens, and 494 MPa for Ti0.7Al0.3Ncoated specimens. Thus the deposition of Ti0.7Al0.3N coating gave an increase of around 10% in fatigue limit. The deposition of the graded (Ti, Al)N coating did not give any increase in fatigue limit.
Fig. 1. Geometry of the rotary bending fatigue test specimen; dimension in mm.
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Table 1 Experimental results for determining the fatigue limit of the uncoated samples. Sample
Alternating stress (MPa)
Number of cycles
1 2 3 4 5 6 7 8 Fatigue limit
450 475 450 475 450 425 450 425 450
11,700,000 7,335,000 11,500,000 255,000 1,740,000 10,000,000 9,500,000 10,000,000
Table 2 Experimental results for determining the fatigue limit of the graded (Ti, Al)N-coated samples. Sample
Alternating stress (MPa)
Number of cycles
1 2 3 4 5 6 7 8 Fatigue limit
450 425 450 475 450 475 450 425 450
4,500,000 10,000,000 10,000,000 9,000,000 10,000,000 4,550,000 8,170,000 10,000,000
Fig. 2. The S–N curves of the substrate, graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens at 500 °C in air. The data used to determine the fatigue limit are indicated by the solid symbols.
Table 4 Parameters involved in the Basquin relationship for the conditions tested.
Table 3 Experimental results for determining the fatigue limit of the Ti0.7Al0.3N-coated samples. Sample
Alternating stress (MPa)
Number of cycles
1 2 3 4 5 6 7 8 Fatigue limit
475 500 525 500 475 500 475 500 494
10,000,000 10,000,000 600,000 9,385,000 10,000,000 3,285,000 10,000,000 575,000
−m
;
A (MPa)
m
741.1 951.4 850.7
0.031 0.046 0.034
3.2. Fractographic analysis
Fig. 2 presents all the data obtained in the fatigue test. Basquin equation, S = ANf
Condition Substrate Graded (Ti, Al)N Ti0.7Al0.3N
After the fatigue test at 500 °C, the steel was slightly oxidized, both on the surface of the specimen and the fracture surface. Fig. 4a illustrates the fracture surface of the uncoated specimen tested at 450 MPa. Fig. 4b is the magnified view of Fig. 4a at the crack initiation site. The photomicrographs show that fracture of the sample occurred due to propagation of a crack which originated from the surface of the sample. A similar conclusion can be drawn from observation of the fracture surfaces of specimens tested at higher stress levels.
ð2Þ
was employed to describe the relationships of the number of cycles to fracture and the alternating stresses of the uncoated and coated specimens, and the curves with internals indicating 95% confidence were plotted in Fig. 2. It can be seen that the fatigue strength was also improved to some extent by the graded coating at higher stress levels (475 MPa–525 MPa). The values of the fatigue strength coefficient A and the fatigue strength exponent m in Basquin relationship were summarized in Table 4. The increase in fatigue life can be computed as: %increase =
Nfcoating −Nfsubstrate Nfsubstrate
× 100%
ð3Þ
where the number of cycles to fracture for each condition is calculated from Eq. (3), employing the values given in Table 4. The results were presented in Fig. 3. For the Ti0.7Al0.3N coating, the increase in fatigue life ranges between 1785 and 2706%, when testing is conducted at stresses between 500 and 575 MPa. For the graded (Ti, Al)N coating, the increase in fatigue life ranges between 112 and 507%, when testing is conducted at stresses between 475 and 525 MPa.
Fig. 3. Change in the percentage of increase in fatigue life for Ti0.7Al0.3N- and graded (Ti, Al)N-coated specimen.
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Fig. 4. (a) SEM fatigue fracture surface and (b) magnified view of the crack initiation site in a for uncoated steel tested at 450 MPa.
Fig. 5. (a) SEM fatigue fracture surface and (b) magnified view of the crack initiation site in a for graded (Ti, Al)N-coated steel tested at 450 MPa.
Although the beneficial effect of the presence of the graded (Ti, Al)N on the fatigue life and strength of the steel is not as significant as that of Ti0.7Al0.3N coating, the fracture surfaces of both kinds of specimens show similar characteristics. The fracture surface appearances of the graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens tested at 450 MPa and 500 MPa are shown in Figs. 5 and 6. It can be seen that for the coated specimen, the fatigue crack propagated from the surface of the substrate into the interior, and eventually caused the failure of the specimen, as shown in
Figs. 5a–b and 6a–b. For the coated specimens tested at higher stress levels, similar characteristics were observed. The coatings remain well adherent to the substrate during the fatigue tests at lower stress level (450 MPa–525 MPa). At higher stress (575 MPa), extensive fracture and spallation of Ti0.7Al0.3N coating occurred. In order to obtain more information about the fatigue crack initiation, surfaces of the graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens, which were tested at 450 MPa and 500 MPa respectively and did not fail over a
Fig. 6. (a) SEM fatigue fracture surface and (b) magnified view of the crack initiation site in a for Ti0.7Al0.3N-coated steel tested 500 MPa.
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Fig. 7. SEM surface morphology of a spalled area on a Ti0.7Al0.3N-coated specimen tested at 500 MPa for 1 × 107 cycles.
1×107 cycle test, were observed. For both kinds of specimens, spallation of the coatings, immediately adjacent to the machining mark, occurred occasionally. The typical view of the spalled areas is shown in Fig. 7. It can be seen that the spallation occurred not along the interface, instead with
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an inclination to it, and the substrate was exposed only at the edge of the groove (machining mark). In some cases, the crack grew within the coatings and led to the spallation, and the substrate was not exposed. The macro-particles which are inevitably deposited during the coating process are also visible, and seem to have negligible effect on the coating spallation and fatigue crack initiation. A fatigue crack, initiated from the substrate surface where the coating had spalled (the edge of the groove), was found in a graded (Ti, Al)N-coated specimen, as shown in Fig. 8a–b. The coating was stripped away by a chemical stripping solution, which did not react with the steel substrate [7], so that more information about the crack can be obtained. As shown in Fig. 8c, the length of the crack is about 25 μm longer than that observed before the coating stripped away, indicating crack propagation in the substrate underneath the coating (Fig. 8d). The cross-sections of specimens were also prepared normal to the surface, taken along the axis of the specimen. Fig. 9a shows cracks within the coating and Fig. 9b shows the cross-section of a typical spalled area shown in Fig. 7. It indicates that the cracks may initiate within the coating or at the coating/substrate interface, and propagate within the coating rather than along the interface, eventually reach the coating surface and lead to the spallation. It also indicates that the bond strength between the coating and the substrate is quite strong. It is worth to mention that, for the graded (Ti, Al)N- and Ti0.7Al0.3N-coated specimens tested at 450 MPa and 500 MPa for 1×107 cycles respectively and did not fail, the spallation of the coating, as shown in Fig. 7, was observed in all the
Fig. 8. (a) SEM surface morphology of a crack originated from the substrate surface where the coating had spalled on a graded (Ti, Al)N-coated specimen tested at 450 MPa for 1 × 107 cycles; (b) magnified view of a; (c) corresponding BSE view of a after coating stripped; and (d) magnified view of the square area in c.
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Fig. 9. BSE cross-sections normal to the specimen surface, taken along the axis of coated specimens. (a) graded (Ti, Al)N-coated specimen tested at 450 MPa for 1 × 107 cycles and (b) Ti0.7Al0.3N-coated specimen tested at 500 MPa for 1 × 107 cycles.
specimens, but the occurrence of fatigue crack, as shown in Fig. 8, was only observed in one specimen. Cross-sections of the failed specimens in the fatigue test were also prepared. Fig. 10 is the cross-section of a Ti0.7Al0.3N-coated specimen tested at 500 MPa and failed at 3.29 × 106 cycles. It can be seen the fatigue crack originated from the substrate surface had propagated into the substrate even when through-scale crack has not formed in the coating well adherent to the substrate. 4. Discussion From Fig. 2, it can be seen the standard deviation of the number of cycles to fracture at various maximum alternating stresses is quite large for uncoated and coated specimens. This may be related with the surface roughness of the specimen. It has been measured that the average roughness at the constant cross-sectional gauge length of the specimens was approximately Ra = 0.35 μm after polishing, but in some areas Ra could reach to 0.67 μm (Fig. 9). After the deposition of the coating, the average surface roughness measured was almost as same as that of the uncoated specimen, and from the cross-sections of the coated specimens (Fig. 9), it can also be observed that the deposition of the coatings did not change the surface roughness obviously. However, most of the experimental data were in or near to the intervals of 95% confidence after S–N curves were plotted employing Basquin relationship. So, from the statistical point of view, it can be included the fatigue life and strength was improved apparently by Ti0.7Al0.3N coating at all the
Fig. 10. BSE cross-section normal to the specimen surface, taken along the axis of Ti0.7Al0.3N-coated specimen tested at 500 MPa and failed at 3.29 × 106 cycles.
testing stress levels, and improved to some extent by the graded coating at higher stress levels (475 MPa–525 MPa). The deposition of Ti0.7Al0.3N coating also gave an increase of around 10% in fatigue limit. It has been measured that the average residual compressive stresses in graded (Ti, Al)N and Ti0.7Al0.3N coatings were around 4 GPa at room temperature [7]. At 500 °C, the compressive stresses in the coatings will be lower than that at room temperature due to the thermal expansion discrepancy between the coating and the substrate, and the stress relaxation (SR) in the coating resulting from the creep deformation of the substrate near the interface at high temperature [17]. Thermal stress can be quantified by the following equation [21]: σth =
Ef ðα −αs ÞðTt −Tr Þ: 1−νf f
ð4Þ
In the above equation, Ef and νf are Young's modulus and Poisson ratio of film, respectively; αf and αs are thermal expansion coefficients of film and substrate; Tt is the temperature of the fatigue test and Tr room temperature. If we use αs = 11.7 × 10− 6 °C− 1, αTiN = 9.4 × 10− 6 °C− 1, αTiAlN = 7.5 × 10− 6 °C− 1, ν = 0.2, and E = 450 GPa [22] in Eq.(4), σth in TiN is about 0.61 GPa and in Ti0.5Al0.5N 1.12 GPa. Suh et al. [17] measured the stress relaxation of TiAlN coating at high temperature and found that the decrement of the residual stress in the coating was around 0.5 GPa after SR treatment at 500 °C. Taking the above two factors into account, the average stresses in the coatings were still compressive at 500 °C. The micro-hardness was 25.1 GPa for the graded coating [19] and 23.5 GPa for Ti0.7Al0.3N coating [20] respectively at room temperature. The difference of the two coatings in micro-hardness was not significant. Cracking and spallation of the coatings occasionally observed on the coated specimens which did not fail in the fatigue test is due to the propagation of the preexisting defects within the coatings or at the coating/substrate interface under the cyclic loading at the sites of stress concentration due to localized change of curvature, for example, at the machining marks. The spallation of the coatings occurred not along the interface instead with an inclination indicated that the coating–substrate interface is quite strong and more resistant to the rupture than is the coating itself. The reason for the cracks in the coatings having not penetrated into the substrate may lie in that the propagation directions of the cracks were not perpendicular to the interface. The fatigue crack was observed to initiate at the substrate surface where the coating had spalled under the cyclic loading on the coated specimen which did not fail in the fatigue test. Furthermore, the fatigue cracks had initiated and propagated into the substrate even
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when the through-scale crack has not formed in the coating which was well adherent to the substrate on the fractured specimens. At 500 °C, the elastic modulus (142 GPa) of 1Cr11Ni2W2MoV steel decreased significantly compared with that at room temperature (196 GPa). The substrate became softer and easier to deform under the alternating stress. The presence of the coatings with high compressive stress will restrict the deformation of the substrate. At the place where the coating had spalled, the restriction to the substrate diminished and the fatigue crack was easier to initiate. At the place where the coating maintained adherent to the substrate, when the plastic deformation of the substrate was too large to be accommodated by the coating, fracture of the coating occurred, and fatigue crack initiated and propagated into the coating and the substrate. The fatigue cracks even had propagated into the substrate when the through-scale crack has not formed in the coating. The compressive coating may retard the growth of surface crack tip to some extent (Fig. 8). Baragetti [1] found that the specified crack propagation depth of CrN-coated steel was deeper than that of the substrate steel and it was the high residual stresses at the surface that retarded the opening of the crack. Tests to clarify the effect of the coatings on crack growth rate in the substrate have been carried out by Kim et al. [2] and the results showed that the presence of the coating layer did not significantly retard crack growth. Therefore it is reasonable to think the presence of the coating will retard the growth of the crack at the very beginning and afterwards its contribution to the restriction of crack growth will be limited. So it can be concluded that main reason the presence of the graded (Ti, Al)N and Ti0.7Al0.3N coatings improving the fatigue life and strength of the steel is that the coatings, which are well adhered and have high residual stress, restricted the deformation of the substrate and delayed the fatigue crack initiation. During fatigue tests at 500 °C, the uncoated specimens were slightly oxidized, which promoted the initiation of the fatigue cracks, but the oxidation of the coatings was very slight, and seems to have negligible effect on the fatigue initiation. This may also account for the improvement of the fatigue properties by the presence of the coatings. The beneficial effect of the presence of the graded (Ti, Al)N on the fatigue life and the fatigue strength of the steel is not as significant as that of Ti0.7Al0.3N coating, which may be related to the difference on the compressive stress in the coatings [7]. For the graded coating, Al content in the coating increased gradually from substrate/coating interface to coating surface, and Ti concentration changed in the opposite direction, thus the inner layer rich in TiN and outer layer rich in Ti0.5Al0.5N. Compared with Ti0.7Al0.3N coating, the stress maximum value in the graded coating was higher and nearer to the coating surface, so the graded coating may be easier to crack and spall under the alternating stress and the restraint to the substrate decreased [7].
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5. Conclusions 1. At 500 °C, the fatigue life and strength of 1Cr11Ni2W2MoV stainless steel were apparently increased by the presence of Ti0.7Al0.3N coating. The fatigue life and strength of the steel was also improved to some extent by the presence of the graded coating at higher stress levels (475 MPa–525 MPa). 2. The presence of the coatings, which were well adhered and have high compressive stress, restricted the plastic deformation of the substrate, thus improved the fatigue properties of the steel.
Acknowledgement This work was supported by the National Nature Foundation of China under grant no. 50671110. The authors greatly appreciate the very helpful discussion with Professors Guangping Zhang, Weicheng Yu, Suhua Ai and Huihe Su about the fatigue results, and wish to thank Mr. Shichen Wang for the coating deposition.
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