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Effect of AL3Ti diffusion aluminide coating on tensile properties of a near α-Ti alloy
Materials Science and Engineering A 530 (2011) 565–573
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Effect of AL3 Ti diffusion aluminide coating on tensile properties of a near ␣-Ti alloy Chandrakant Parlikar ∗ , Md Zafir Alam, Dipak K. Das Defence Metallurgical Research Laboratory, P.O.-Kanchanbagh, Hyderabad 500058, India
a r t i c l e
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Article history: Received 3 February 2011 Received in revised form 7 October 2011 Accepted 10 October 2011 Available online 18 October 2011 Keywords: Near ␣ titanium alloy Al3 Ti coating Tensile properties Oxidation
a b s t r a c t The tensile properties of a near ␣ Ti-alloy have been evaluated in both uncoated and diffusion aluminide (Al3 Ti) coated conditions at room temperature (RT) and 600 ◦ C. The effect of coating thickness on the tensile properties has been examined. The YS and UTS of the alloy at RT reduced marginally because of the presence of the coating, with the extent of reduction increasing with the increase of coating thickness. Such a trend in alloy strength was also observed at 600 ◦ C. The RT ductility of the alloy remained nearly unaffected by the presence of the coating. However, an increase in ductility was observed in the coated alloy at 600 ◦ C. The above effect of Al3 Ti coating on tensile properties of the present near ␣ Ti-alloy has been explained in terms of the influence of the through-thickness cracking in the coating during tensile testing. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Near ␣-Ti alloys such as IMI-834 and Timetal 1100 are widely used for the fabrication of compressor parts of advanced gas turbine engines. However, the use of these alloys is limited to a temperature of about 600 ◦ C because of their inherently poor oxidation resistance [1]. Oxidation causes the formation of a brittle ␣ casing which adversely affects the mechanical properties of these alloys [2]. Increasing the high temperature use of near ␣-Ti alloys is being attempted by application of oxidation resistant coatings. Several coatings such as TiAlN, Al3 Ti type aluminides, silicides and MCrAlY type coatings have been reported for this purpose [3–14]. Al3 Tibased diffusion aluminide coatings formed via pack aluminizing method have been reported to greatly enhance the oxidation resistance of IMI-834 alloy up to 800 ◦ C [13,14]. Both unmodified and Pt-modified Al3 Ti coatings have been evaluated for their oxidation resistance on the above alloy and have been found to provide excellent high temperature protection to the alloy [13,14]. These coatings derive their oxidation resistance from their ability to form an adherent layer of Al2 O3 on the surface, which, once formed, retards further oxidation [14–16]. Despite their good oxidation resistance, Al3 Ti coatings are inherently brittle and develop through-thickness cracks during coating formation as well as during cyclic oxidation exposure [13,17]. Therefore, these coatings can potentially degrade the mechanical properties of the substrate alloy. Examples of brittle diffusion
∗ Corresponding author. Tel.: +91 40 24586839. E-mail address: [email protected] (C. Parlikar). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.10.021
aluminide coatings such as -NiAl degrading the mechanical properties of substrate Ni-base superalloys can be found in the reported literature [18–20]. The objective of the present study is to evaluate the effect of Al3 Ti diffusion aluminide coating on the tensile properties of a near ␣-Ti alloy at room temperature (RT) and 600 ◦ C. The effect of coating thickness on the tensile properties has also been examined in this study. 2. Experimental details An experimental near ␣ Ti-alloy having a nominal composition (in wt.%) as Ti–5.93Al–4.08Sn–3.93Zr–0.71Nb– 0.59Mo–0.37Si–0.076C, was used as the substrate material for the present study. The dissolved oxygen and nitrogen contents in the alloy were 826 and 32 ppm, respectively. The extruded alloy rods (23 mm diameter) were given the required heat treatment, i.e. solutionized at 1000 ◦ C for 2 h followed by aging at 700 ◦ C for 2 h. After both these treatments, the rods were cooled outside the furnace to room temperature in ambient air. The solutionizing and aging treatments were also carried out in air. The volume fraction of equiaxed primary ␣ in the heat treated alloy was approximately 0.3. Similar near ␣ alloys having a relatively high volume fraction of equiaxed primary ␣ phase (0.25–0.3) have been employed to achieve a good combination of creep and fatigue properties [21,22]. Cylindrical blanks of 75 mm length and 8 mm diameter were removed from the heat treated rods using electro-discharge machine (EDM) technique. Round tensile specimens having a gage length of 28 mm, parallel length of 20 mm and gage diameter of 4 mm, as per ASTM standard E8M [23], were then fabricated from the blanks in a computerized numerical-controlled (CNC) lathe
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Fig. 1. Cross-sectional microstructure of the Al3 Ti coating formed at an aluminizing temperature of 700 ◦ C. Through-thickness cracks in the coating can be seen.
machine. Some of the tensile specimens were applied with an aluminide coating by pack aluminizing method. The powder pack used for aluminization consisted of (in wt.%) 15% Al as the Al source, 2% NH4 Cl as the activator and 83% calcined Al2 O3 as the filler. The pack aluminizing process was carried out in a tube furnace. In order to obtain coatings of different thicknesses, aluminization was carried out at three different temperatures, namely 600, 650 and 700 ◦ C. The duration of aluminizing was kept constant at 5 h in all the cases. To prevent oxidation during the process, an Ar atmosphere was maintained in the tube furnace. The coated specimens, along with the uncoated ones, were tensile tested at room temperature (RT) and 600 ◦ C by using a universal testing machine (model Instron 5500R). All the tensile tests were carried out at an initial strain rate of 6.7 × 10−4 s−1 . For testing at 600 ◦ C, the tensile specimen was first fitted to the grips in the machine and then surrounded by a split-type furnace. The temperature of the furnace was then raised from RT to 600 ◦ C, which took about 1 h. During the heating-up period, the sample remained inside the furnace. After the furnace reached the above temperature, the sample was soaked at this temperature for approximately 15 min to attain temperature uniformity. Subsequently, it was tensile tested in a similar manner as mentioned earlier. At least two samples were tested under any given condition for ensuring consistency in the results. Coating microstructure was examined
Fig. 2. Composition of Al and Ti across the thickness of the coating formed at 700 ◦ C.
Fig. 3. Tensile properties of the uncoated and coated alloy: (a) at RT and (b) at 600 ◦ C.
with the help of a Quanta 400 scanning electron microscope (SEM) in as-coated condition as well as after tensile test. The phase constitution of the coating was determined by X-ray diffraction (XRD) method. A Cameca SX-100 electron probe micro-analyzer (EPMA) was utilized to analyze the composition of the coating. To examine surface cracking, the gage length portion of the failed tensile samples was observed in the above SEM. The fracture surfaces of the failed samples were also examined in the SEM.
Fig. 4. Engineering stress–strain plots for UC and C35 samples at both RT and 600 ◦ C.
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Table 1 Details of all the tensile test specimens used in the present study. Sample ID
Sample details
UC C5 C18 C35 UC5 UC18 UC35
Uncoated samples prepared from fully heat treated rods Samples with 5 m thick coating Samples with 18 m thick coating Samples with 35 m thick coating Tensile specimen prepared after machining out a 0.2 mm thick surface layer from the parallel length of a C5 sample Tensile specimen prepared after machining out a 0.2 mm thick surface layer from the parallel length of a C18 sample Tensile specimen prepared after machining out a 0.2 mm thick surface layer from the parallel length of a C35 sample
3. Results 3.1. As-coated microstructure The cross-sectional microstructure of the as-deposited coating corresponding to aluminization temperature of 700 ◦ C is shown in Fig. 1. The coating was about 35 m thick and had a single layer consisting of Al3 Ti phase, as determined from the XRD analysis [24,25]. The above phase constitution of the coating was also confirmed by EPMA chemical analysis, as shown in Fig. 2. The atom ratio of Al to Ti as 3:1 indicates the coating phase to be Al3 Ti. The coatings formed at 600 and 650 ◦ C also had the identical single-layer microstructure and consisted of Al3 Ti phase, although their thicknesses were lower at 5 and 18 m, respectively. Increase in coating thickness achieved by increasing the aluminizing temperature was consistent with the corresponding increase of Al intake by the sample during coating process [24,25]. The Al pick-up registered by the tensile specimens for the above three temperatures was 1.42, 4.63 and 8.1 mg cm−2 , respectively. Several through-thickness cracks were present in the as-deposited coatings (Fig. 1). These cracks had
formed due to the mismatch in the coefficient of thermal expansion (CTE) between the Al3 Ti coating (CTE = 15 × 10−6 ◦ C−1 [26]) and the Ti-alloy substrate (CTE = 8.9 × 10−6 ◦ C−1 [27]). Tensile stresses that were generated in the coating due to the above CTE mismatch during cooling of the coated sample after the coating process, caused the above mentioned through-thickness cracking in the brittle Al3 Ti coating [13,17,24,25]. 3.2. Tensile properties The tensile properties of the uncoated and coated alloy were determined from the engineering stress–strain plots. The properties at RT are presented in Fig. 3(a). While the uncoated tensile specimens have been referred to as UC in this study, the coated ones with coating thicknesses of 5, 18 and 35 m have been referred to as C5, C18 and C35, respectively, as indicated in Table 1. The 0.2% yield strength (YS) of the uncoated alloy was 1098 MPa and that for C5 was very similar at 1106 MPa. The YS of C18 was, however, 3% lower than that for UC alloy and the same for C35 was lower by about 12%. Thus, it is evident that, with increase in coating
Fig. 5. Slip patterns generated on the surface of the uncoated samples during the tensile test: (a) at RT (low magnification image), (b) at RT (high magnification image) and (c) at 600 ◦ C. Arrows indicate cracks in (c).
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Fig. 6. Surface morphology of the coated samples in the gage length portion: (a) C5 prior to tensile test at RT, (b) C35 after test at RT, (c) C35 after testing at 600 ◦ C. Arrows indicate cracks in (a) and (b).
thickness, the YS of the coated alloy progressively decreased. Similarly, a progressive decrease in the ultimate tensile strength (UTS) of the alloy, although marginal (<5%), was also observed as the coating thickness was increased. This is evident from the UTS values of UC, C5, C18 and C35 which were 1194, 1175, 1159 and 1141 MPa, respectively (see Fig. 3(a)). Despite causing a decrease in strength (both YS and UTS), the presence of the coating did not cause any appreciable decrease in the plastic strain to fracture of the alloy (henceforth referred to as ductility) in case of C5 and C18, as the ductility values of about 10% for these samples were similar to that of UC. However, the ductility of C35 was somewhat lower at 8% as compared to the value of 12% measured for UC (Fig. 3(a)). It has been reported that precipitation of ordered Ti3 Al (␣2 ) phase and silicides occurs in near ␣ Ti-alloys during the aging treatment [22,28,29]. The precipitated ␣2 and silicide particles and their coarsening are believed to cause some degree of increase in the tensile strength and reduction in the ductility of the alloy both at RT and high temperatures [22,30]. The present substrate alloy, in case of the coated samples, was subjected to the treatment of 5 h at 600/650/700 ◦ C during the coating formation process. This treatment was in addition to the aging treatment (2 h at 700 ◦ C) given to the alloy prior to coating formation. To examine whether the possible changes in the alloy microstructure (in terms of precipitation/coarsening of ␣2 and silicide particles [28,31]) because of the above additional exposure during coating formation has caused any change in the tensile properties of the alloy, the coating was completely removed from a set of coated tensile samples by machining out a 0.2 mm thick layer from the surface over the entire parallel length. These samples have been referred to as UC5, UC18 and
UC35, signifying the initial thickness of the coating that was present on these samples before its removal (Table 1). These samples were then tested at RT in a similar manner as described earlier for C5, C18 and C35. The tensile properties of these samples are also shown in Fig. 3(a). It is clear that strength values (both YS and UTS) of UC5, UC18 and UC35 were very similar to each other (within ±5%) and to those of UC. The ductility values of UC5, UC18, UC35 were also very similar to that of UC and lied in the range 10–13%. This observation clearly indicates that the high temperature exposure during the coating treatment, i.e. 5 h at 600/650/700 ◦ C, did not cause any appreciable change in the tensile properties of the substrate alloy. Therefore, the minor degradation in the tensile properties of the alloy observed in the coated condition, i.e. in samples C5, C18 and C35, as mentioned previously, was primarily caused by the presence of the coating. Similarly, the somewhat lowering of ductility in case of C35, as mentioned earlier, was also caused by the coating, although such an effect was not prominent in case of C5 and C18. The tensile properties of the present alloy at 600 ◦ C are shown in Fig. 3(b). The YS and the UTS for the uncoated (UC) alloy at the above temperature were 597 and 721 MPa, respectively. Similar to that observed at RT, a reduction in alloy strength (YS and UTS) was also noticed at 600 ◦ C. The trend of progressive decrease in YS with increase in coating thickness was also observed at 600 ◦ C, as evident from the YS values of 580, 567 and 523 MPa measured for C5, C18 and C35, respectively (Fig. 3(b)). A similar trend was also maintained in the UTS values of the coated alloy (Fig. 3(b)), with UTS of C35 at 590 MPa, being nearly 20% lower than that of UC. It is evident from Fig. 3(b) that the ductility value of 11.5% measured for UC at 600 ◦ C was very similar to that measured at
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RT (Fig. 3(a)). However, the coated alloy showed a much higher ductility at 600 ◦ C than at RT. The ductility of the coated alloy at 600 ◦ C lied in the range 14.5–20.5%, which was significantly higher than the corresponding value of 11.5% for UC. Further, among the coated samples, a decrease in the ductility value from about 20% to 14.5% was observed as the coating thickness increased from 5 m to 35 m. As was done at RT, the coating-removed samples UC5, UC18 and UC35 samples were also tested at 600 ◦ C and the tensile properties of these samples are shown in Fig. 3(b). As clearly evident, YS, UTS and ductility of UC5, UC18 and UC35 were very similar to the corresponding values of UC. This again suggests that the degradation in strength and the enhancement in ductility observed for the coated alloy at 600 ◦ C were primarily caused by the presence of the coating. The typical tensile curves for uncoated and coated samples corresponding to RT and 600 ◦ C are shown in Fig. 4. The curves for the coated sample in the above figure are for C35. As typical of near ␣ Ti-alloys [32], the RT tensile curve for UC did not show any appreciable sign of post-UTS necking till failure. The behavior of the RT tensile plot for C35 sample was also very similar to that of UC. Both uncoated and the coated samples exhibited enhanced post-UTS necking when tested at 600 ◦ C, as evident from decreasing stress with increasing strain beyond UTS in Fig. 4. The tensile behavior of the present near ␣ Ti alloy in uncoated and coated conditions can be summarized as follows. Application of Al3 Ti coating on the alloy led to a marginal (up to 12%) reduction in strength (both YS and UTS) at RT, with the extent of reduction increasing with the increase in coating thickness. The presence of the coating did not cause any appreciable reduction in alloy ductility at RT for C5 and C18, which remained in the range of 10–12%. Only in the case of C35, the ductility reduced to 8% as compared to the value of 12% for the uncoated alloy. The variation of strength of the coated alloy at 600 ◦ C was similar to that observed at RT. However, unlike at RT, the presence of coating led to an enhancement in the ductility of the alloy at 600 ◦ C, although a trend of decrease in ductility with increase in coating thickness was observed at this temperature. 3.3. Microstructural and fractographic observations in tensile samples 3.3.1. Surface features over gage length Typical surface features on the gage length portion of the failed UC samples are shown in Fig. 5(a). Numerous near-parallel lines were found criss-crossing the surface of the tested samples. These marks were not pre-existing on the as-machined samples and were generated during the tensile test. A magnified view of the surface, as shown in Fig. 5(b), suggests the above marks to be sets of parallel slip lines distributed over the sample surface. The angles between the intersecting sets of parallel slip lines were measured to be 53 −
−
and 55◦ . It has been reported that usually {0 0 0 1}, {1 0 1 0}, {1 0 1 1} −
and {1 1 2 2} slip planes operate during deformation in ˛/ˇ Ti alloys [33]. The interplanar angle () between these slip planes can be calculated by using the following equation (Eq. (1)) [34]: 2
hd + ke + (1/2)(he + kd) + (3/4)lg(a/c)
cos =
2 1/2
{h2 + k2 + hk + (3/4)l2 (a/c) }
2 1/2
(1)
× {d2 + e2 + de + (3/4)g 2 (a/c) }
The indices (hkil) and (defg) in the above equation denote the Miller indices for two planes between which the angle is being calculated. The measured angles, as mentioned above, corresponded to comparison of the above measured values of 53 and 55◦ with various interplanar angles indicated that the participating slip planes in the −
deformation of UC sample during tensile test were {0 0 0 1}, {0 1 1 1} −
Fig. 7. Longitudinal sections of the tensile specimens tested at RT: (a) UC, (b) C5 and (c) C35. Magnified view showing voids is presented in the inset in (a).
and {1 0 1 0} type. Same slip line pattern was also observed on the UC sample tested at 600 ◦ C, as evident from Fig. 5(c). Additionally,
this sample showed extensive surface cracking with the cracks oriented transverse (perpendicular) to the loading direction, as shown in the above figure. Cracking along slip lines was also observed at several locations. A few fine cracks could be located on the surface of as-coated specimens (C5, C18 and C35), as typically shown in Fig. 6(a) for C5 sample. These cracks correspond to the through-thickness cracks generated in the coating under the influence of CTE mismatch, as
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Fig. 8. Longitudinal sections of the tensile specimens tested at 600 ◦ C: (a) UC, (b) C5 and (c) C35.
mentioned earlier (Fig. 1). During tensile test at RT, many additional cracks were generated in the coating transverse to the loading direction, as typically shown in Fig. 6(b) for C35. The cracks generated during RT testing largely remained fine in all the three coated samples. No slip traces, as seen in UC sample, were observed on the surface of the failed coated samples. The coated samples tested at 600 ◦ C also revealed large number of transverse cracks on the surface, as typically shown in Fig. 6(c) for failed C35 sample. These cracks were also much wider than those generated by testing at RT (Fig. 6(b)). 3.3.2. Observations on longitudinal-sections The longitudinal sections of the failed UC, C5 and C35 samples, tested at RT, are presented in Fig. 7(a)–(c). It is evident from Fig. 7(a) that no cracking developed at the surface, either close to fracture end or away from it, prior to the failure of UC sample. This is unlike the case in C5, C18 and C35 where many through-thickness cracks of the coating could be found penetrating into the substrate over the entire gage length, as typically shown in Fig. 7(b) and (c) for C5 and C35, respectively. These through-thickness cracks could also be seen on the sample surface, as mentioned previously (Fig. 7(b)). In case of the uncoated sample, several voids were found to be present in the substrate close to the fracture end (Fig. 7(a)). A closer examination revealed that the voids had formed by de-cohesion of the ␣ lath boundaries within the transformed  colonies. Fig. 8 presents the longitudinal section of a failed UC sample along with those of C5 and C35 samples tested at 600 ◦ C. Many surface cracks, penetrating into the substrate could be observed, throughout the longitudinal section, in case of UC (Fig. 8(a)). These cracks corresponded to the transverse cracks that were seen on the surface, as mentioned earlier (Fig. 5(c)). The cracks seen in the longitudinal section of UC, largely remained at the surface and did not
penetrate to any appreciable depth into the sample, except close to the fracture end. In case of the three coated samples, throughthickness cracks were present in the coating all along the gage length and some coating cracks were found to have penetrated into the substrate, especially in C35 (Fig. 8(c)). Apart from the cracks at the surface, voiding in the substrate close to the fracture end was observed in both uncoated and coated samples that were tested at 600 ◦ C. In order to understand the effect of cracks in the coating on the strength of the alloy, especially UTS, a few tests were carried out at RT and 600 ◦ C wherein the loading was stopped very close to UTS. Subsequently, the samples were unloaded and their longitudinal sections metallographically polished and inspected in SEM. A large number of through-thickness cracks along the gage length could be found in all the coatings at both the temperatures, similar to that mentioned earlier for the failed specimens (Fig. 9(a)). Further, the penetration of the cracks into the substrate was also observed in all the three cases at both the temperatures. However, at 600 ◦ C, the number of penetrating cracks and their penetration depths into the substrate were much higher for C35 than for either C5 or C18, as evident in Fig. 9(b). In fact, most of the cracks in the latter two cases remained restricted to within the coating (Fig. 9(a)). 3.3.3. Fractography Fig. 10(a) presents the fracture surface of a failed UC sample tested at RT. The presence of quasi-cleavage features over the entire fracture surface was observed in this sample. The fracture surface of all the coated specimens tested at RT, including C35, also exhibited similar quasi-cleavage features, as typically shown in Fig. 10(b). The separation along the interfaces of ␣ laths in the transformed  colonies could be clearly seen on the cleavage planes, as shown in Fig. 10. Unlike at RT, the failure at 600 ◦ C
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Fig. 9. Longitudinal sections of the tensile specimens after interrupted tensile testing at 600 ◦ C: (a) C5 and (b) C35.
occurred in a ductile manner in all the four cases (UC and the three coated samples). This was evident from the dimpled appearance of the fracture surfaces for these samples, as typically shown in Fig. 11(a) for UC and Fig. 11(b) for C35. The presence of several voids in the substrate close to the fracture end in all the failed samples, as mentioned earlier, supports the above ductile failure mechanism involving the nucleation of voids and their subsequent coalescence causing fracture. 4. Discussion As mentioned previously, the substrate alloy underwent an additional exposure for 5 h at the aluminizing temperature (600/650/700 ◦ C), which was close or equal to its aging temperature. It has been reported that, in near ␣ alloys, some degree of precipitation and subsequent coarsening of Ti-silicides and ␣2 (Ti3 Al) phase occur during the aging treatment as well as subsequent use at high temperatures [22,28,29]. Zhang et al., in case of two near ␣ alloys, have shown that such precipitation and/or coarsening during the aging treatment did not cause any appreciable change in alloy strength and ductility, irrespective of the duration of aging (which was varied in the range 2–15 h at aging temperatures of 700 and 760 ◦ C) [22]. Only when long term thermal exposure at 600 ◦ C for 100 h was given, an increase in alloy strength and a decrease in ductility was observed, which was ascribed to additional precipitation/coarsening of silicides and ␣2 particles. Woodfield et al. have also made similar observations in case of near ␣ alloy Ti 5331S [30]. Based on these reported studies, it is clear that
Fig. 10. Fracture surfaces of the failed tensile samples tested at RT: (a) UC and (b) C35. Cleavage separation along the ␣ lath boundaries within transformed  colonies can be seen from the inset in Fig. 9(b).
the effect of aging temperature and its duration on the strength and ductility of near ␣ alloys is at best limited. Therefore, it can be assumed that the additional high temperature exposure of the present substrate alloy during the aluminizing treatment would not cause any appreciable change in its tensile properties. This aspect is also clearly evident from the fact that the strength and ductility values of UC, UC5, UC18 and UC35 were very similar at both RT and 600 ◦ C (Fig. 3). Being a diffusion coating, the present Al3 Ti coating formed an integral part of the substrate. As mentioned earlier, the brittle coating, irrespective of its thickness, had several through-thickness cracks in as-formed condition (Fig. 1). Since Al3 Ti phase remains brittle up to 800 ◦ C [35,36], additional cracks were generated in the coating during very early stages of tensile loading at both the test temperatures (Figs. 7 and 8). Because of the presence of the through-thickness cracks, the coating in the coating-substrate composite structure, i.e. in the coated samples, can be considered as non-load bearing during the tensile test. However, the coating cracks would cause stress concentration at the crack tips, i.e. in the coating/substrate interface. Because of stress concentrations at the crack tips, the substrate would locally yield at a stress lower than
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Fig. 11. Fracture surfaces of the failed tensile samples tested at 600 ◦ C: (a) UC and (b) C35.
the YS of the uncoated alloy. Thus, the YS of the alloy would decrease in coated condition, as was observed in the present study (Fig. 3). The extent of YS drop in the coated samples would depend on the magnitude of the stress concentration at the crack tips, which is function of the crack length (or coating thickness) and crack tip radius. Assuming the crack tip radius to be equal in case of all the three coated samples in as-formed condition, the stress concentration at the crack tips would have a direct dependence on the crack length [37]. Thus, the decreasing trend in the YS observed in the coated samples with increase in coating thickness was due to the corresponding increase in the stress concentration ahead of the through-thickness crack tips present in the coatings. There are two major factors that are likely to cause a reduction in the UTS values of the coated samples: (1) the non-load bearing nature of the coating and (2) the penetration of the coating cracks into the substrate leading to further lowering of the load bearing cross-section of the sample. Thus, it can be concluded that the reduction in the UTS values of the coated samples observed in this study (with respect to the UTS of UC) has been a combined contribution of the above two factors. To evaluate the contribution of the non-load bearing coating, the coating thickness (t) can
be subtracted from the initial diameter (D) of the coated sample and the UTS can be revised upward based on the reduced diameter, i.e. D − 2 × t. The revised UTS values for C5, C18 and C35 at RT are 1181, 1180 and 1182 MPa, respectively. These values are very similar (within ±2%) to 1194 MPa measured for the UC sample (see Fig. 3(a)). Thus, it is clear that the presence of the non-load bearing coating was primarily responsible for causing a reduction in the UTS of the coated samples at RT, with the degree of reduction directly dependent on the coating thickness. The contribution of the second factor, i.e. the lowering of the load bearing cross-section of the substrate because of the penetration of coating cracks, to the above reduction in UTS was minimal. The UTS values for the coated samples at 600 ◦ C can also be revised upward by subtracting the coating thickness from the sample diameter, as mentioned above. The revised UTS values at 600 ◦ C for C5, C18 and C35 are 651, 649 and 611 MPa, respectively. These revised values, which include only the contribution of the nonload bearing nature of the coating, are 10–15% lower than the UTS value of 721 MPa measured for UC sample (Fig. 3(b)). This difference between the revised values and the UTS of uncoated sample has actually been contributed by the lowering of the load bearing cross-section of the substrate because of the penetration of coating cracks. The decreasing trend in the UTS of coated samples at 600 ◦ C with increasing coating thickness was in agreement with the previously mentioned observation that the crack penetration depth into the substrate increased with coating thickness during the tensile test (Fig. 8). The effect of through-thickness cracks in the coating and their penetration into the substrate, as mentioned above, did not cause any significant reduction in the ductility in C5 and C18 specimens at RT (Fig. 3(a)). However, the somewhat lower ductility of C35 as compared to UC was possibly because of the much higher crack length and the corresponding higher stress concentration ahead of the cracks in the coating. From Fig. 10(a), it appears that the decrease in ductility in case of C35 was too small to be reflected in terms of any change in fracture surface features. The uncoated alloy was expected to show a comparatively higher ductility at 600 ◦ C than at RT. However, the ductility of UC at both temperatures was found to be virtually the same (Fig. 3). As stated previously, for testing at 600 ◦ C, the tensile specimen remained inside the furnace for about 1.25 h during which the furnace attained the above temperature and homogenization of temperature occurred inside the sample. It is believed that, during this period, the surface of the sample underwent some degree of oxidation-related damage, possibly in terms of the formation of a thin ␣ case. Although the presence of an ␣ case could not be confirmed, the several surface cracks observed in the failed sample, as seen in Fig. 5(c) and Fig. 8(a), supports the possibility of the above surface damage to the UC alloy. Thus, even though the alloy inherently has a higher ductility at 600 ◦ C, it was most likely negated by the loss in ductility caused by the above surface damage [2,21]. Consequently, no significant increase in ductility was registered for UC at 600 ◦ C as compared that at RT. In case of the coated specimens, the presence of the protective Al3 Ti coating prevented the substrate alloy to undergo any surface damage during the heating-up period prior to tensile testing. As a result, much higher ductility was obtained in these samples at 600 ◦ C as compared to that in UC. Such increased ductility values of the coated samples are expected to be close to the actual ductility of the substrate alloy at 600 ◦ C, in the absence of surface damage caused by oxidation. It is interesting to note that such increased ductility for the coated alloy was observed at 600 ◦ C despite the presence of numerous through-thickness cracks in all the three coatings (Fig. 6(c)). Increase in the coating thickness and, hence, the length of the through-thickness cracks (Fig. 3(b)) did cause some drop in the ductility of the alloy, as evident from Fig. 3(b). However, the enhancement in ductility resulting from the surface protection
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provided by the coating was much larger to negate the effect of the coating cracks. As a result, an appreciable increase in the overall ductility was registered by the coated alloy at 600 ◦ C for all three coating thicknesses (Fig. 3(b)). The beneficial effect of the coating, as described above, has also been observed by Leyens et al. [38] in their creep study (at 600 ◦ C) on uncoated and aluminide coated near ␣ alloy Timetal 1100. From the present study, it also appears that the cracking in the coating does not significantly affect the ductility of the alloy even at RT. The results of the present study suggests that the application of the diffusion aluminide (Al3 Ti) coating does not cause any significant loss in ductility of the experimental near ␣ Ti-alloy both at RT and 600 ◦ C. However, it leads to a marginal (10–18%) decrease in the alloy strength. From the tensile results of this investigation and the reported excellent oxidation resistance of Al3 Ti coating [13,14], it can be concluded that the above coating shows good promise for providing high temperature oxidation protection to near ␣ Ti-alloys, without causing appreciable degradation in their tensile properties. 5. Conclusions The tensile properties of a near ␣ Ti-alloy have been evaluated in uncoated as well as diffusion aluminide (Al3 Ti) coated conditions at RT and 600 ◦ C. The effect of coating thickness (5, 18 and 35 m) on the tensile properties of the alloy has also been examined. The presence of the coating was found to marginally reduce (by up to 12%) the YS and UTS of the alloy at RT, with the extent of reduction increasing with the increase of coating thickness. Such a trend of decrease in alloy strength with increase in coating thickness was also observed at 600 ◦ C. The effect of coating thickness on RT ductility of the alloy was marginal. Only in the case of highest coating thickness (35 m), the RT ductility reduced to 8% from the value of 12% for the uncoated alloy. However, an enhancement in ductility of the alloy was observed in presence of the coating at 600 ◦ C for all the three coating thicknesses. The drop in YS observed in the coated alloy was explained in terms of the stress concentration caused by the through-thickness cracks present in the coating. The drop in UTS, on the other hand, could be ascribed to the combined effect of non-load bearing nature of the coating and the loss in load bearing cross-section because of the penetration of coating cracks into the substrate. Acknowledgements The authors acknowledge the assistance provided by MBG, SFAG and EMG groups of DMRL. They are thankful to Director, DMRL, for his permission to publish the present work. This research work has been funded by the Defence Research and Development Organization, DRDO.
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References [1] C. Leyens, in: C. Leyens, M. Peters (Eds.), Titanium and Titanium alloys, WileyVCH, Weinheim, 2003, p. 207. [2] W. Jia, W. Zeng, X. Zhang, Y. Zhou, J. Liu, Q. Wang, J. Mater. Sci. (2010), doi:10.1007/s10853-010-4926-1. [3] H. Benien, J. Maushart, M. Meyer, R. Suchentrunk, Mater. Sci. Eng. A 139 (1999) 126. [4] B.G. MacMordie, Surf. Coat. Technol. 49 (1991) 18. [5] A. Takie, A. Ishida, High Temperature Corrosion of Advanced Materials and Protective Coatings, Elsevier Science Publishers B.V., Amsterdam, 1992, p. 317. [6] J. Karc, M. Ferdinand, D. Liska, S. Tulega, High Temperature Intermetallics, The Institute of Metals, London, 1991, p. 191. [7] B. Cockeram, R.A. Rapp, Metall. Mater. Trans. A 26A (1995) 777. [8] B. Cockeram, R.A. Rapp, Oxid. Met. 45 (1996) 427. [9] J.R. Nicholls, M.R. Winstone, M.J. Deakin, S. Kerry, Elevated Temperature Coatings: Science and Technology II, Minerals, Metals and Materials Society/AIME, Warrendale, 1996, p. 199. [10] D.K. Dewald, D.E. Mikkola, Elevated Temperature Coatings: Science and Technology II, Minerals, Metals and Materials Society/AIME, Warrendale, 1996, p. 255. [11] C. Leyens, M. Schmidt, W.A. Kaysser, Mater. Sci. Eng. A 239–240 (1997) 680. [12] C. Leyens, M. Peters, W.A. Kaysser, Titanium’99: Science and Technology, CRISM “Prometey”, vol. II, 2000, p. 866. [13] M.Z. Alam, D.K. Das, Corros. Sci. 51 (2009) 1405. [14] I. Gurrappa, A.K. Gogia, Mater. Sci. Technol. 17 (2001) 581. [15] D.K. Dewald, D.E. Mikkola, Proceedings of Conference on Elevated Temperature Coatings: Science and Technology II, Anaheim, California, USA, Minerals, Metals and Materials Society/AIME, Warrendale, PA, 4–8 February, 1996, pp. 255–264. [16] J.R. Nicholls, M.R. Winstone, M.J. Deakin, S. Kerry, Proceedings of Conference on Elevated Temperature Coatings: Science and Technology II, Anaheim, California, USA, Minerals, Metals and Materials Society/AIME, Warrendale, PA, 4–8 February, 1996, pp. 199–208. [17] S.-C. Kung, Oxid. Met. 34 (1990) 217. [18] A. Strang, E. Lang, in: R. Brunetand (Ed.), High Temperature Alloys for Gas Turbines, D. Reidel Publishing Company, Holland, 1982, p. 469. [19] K. Schneider, H.W. Grunling, Thin Solid Films 107 (1983) 395. [20] J.M. Veys, R. Mervel, Mater. Sci. Eng. 88 (1987) 253. [21] D.F. Neal, Proc. 6th World Conf. on Titanium, vol. 1, Les Editions de Physique, Les Ulis Cedex, France, 1988, pp. 253–258. [22] J. Zhang, N. Peng, Q. Wang, X. Wang, J. Mater. Sci. Technol. 25 (2009) 454. [23] Standard Test Methods for Tension Testing of Metallic Materials, ASTM International, PA, pp. 1–25, 2008 (source: www.astm.org) (accessed 7.08.08). [24] D.K. Das, S.P. Trivedi, Mater. Sci. Eng. A 367 (2004) 225. [25] S.P. Trivedi, D.K. Das, Intermetallics 13 (2005) 1122. [26] M. Yamaguchi, H. Inui, Intermetallic Compounds, Principle Practice, vol. 2, John Wiley and Sons, 1995, pp. 147–173. [27] M.J. Donachie Jr., Titanium A Technical Guide, ASM International Metals Park, OH 44073, 1989, pp. 9–19. [28] K.V.S. Srinadh, N. Singh, V. Singh, Bull. Mater. Sci. 30 (2007) 595. [29] A. Madson, H. Ghonem, Mater. Sci. Eng. A 177 (1994) 63. [30] A.P. Woodfield, P.J. Postans, M.H. Loretto, R.E. Smallman, Acta Metall. 36 (1988) 507. [31] A.K. Singh, T. Roy, C. Ramachandra, Metall. Trans. A 27 (1996) 1167. [32] N. Singh, V. Singh, Mater. Sci. Eng. A 485 (2008) 130. [33] D. Deka, D.S. Joseph, S. Ghosh, M.J. Mills, Metall. Mater. Trans. A 37A (2006) 1371. [34] J.W. Edington, Practical Electron Microscopy in Materials Science, Van Nostrand Reinhold Company, New York, 1985, p. 294. [35] G. Itoh, K. Kita, M. Kanno, in: P.A. Blenkinsop, W.J. Evans, H.M. Flower (Eds.), Titanium’95, vol. I, The Institute of Materials, London, 1996, pp. 581–587. [36] M. Yamaguchi, Y. Umakoshi, Prog. Mater. Sci. 34 (1990) 1. [37] N.A. Noda, M. Sera, Y. Takase, Int. J. Fatigue 17 (1995) 163. [38] C. Leyens, M. Peters, W.A. Kaysser, Proc. 8th world Conf. on Titanium, Birmingham, UK, 1995, pp. 1935–1942.
Contents lists available at SciVerse ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of AL3 Ti diffusion aluminide coating on tensile properties of a near ␣-Ti alloy Chandrakant Parlikar ∗ , Md Zafir Alam, Dipak K. Das Defence Metallurgical Research Laboratory, P.O.-Kanchanbagh, Hyderabad 500058, India
a r t i c l e
i n f o
Article history: Received 3 February 2011 Received in revised form 7 October 2011 Accepted 10 October 2011 Available online 18 October 2011 Keywords: Near ␣ titanium alloy Al3 Ti coating Tensile properties Oxidation
a b s t r a c t The tensile properties of a near ␣ Ti-alloy have been evaluated in both uncoated and diffusion aluminide (Al3 Ti) coated conditions at room temperature (RT) and 600 ◦ C. The effect of coating thickness on the tensile properties has been examined. The YS and UTS of the alloy at RT reduced marginally because of the presence of the coating, with the extent of reduction increasing with the increase of coating thickness. Such a trend in alloy strength was also observed at 600 ◦ C. The RT ductility of the alloy remained nearly unaffected by the presence of the coating. However, an increase in ductility was observed in the coated alloy at 600 ◦ C. The above effect of Al3 Ti coating on tensile properties of the present near ␣ Ti-alloy has been explained in terms of the influence of the through-thickness cracking in the coating during tensile testing. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Near ␣-Ti alloys such as IMI-834 and Timetal 1100 are widely used for the fabrication of compressor parts of advanced gas turbine engines. However, the use of these alloys is limited to a temperature of about 600 ◦ C because of their inherently poor oxidation resistance [1]. Oxidation causes the formation of a brittle ␣ casing which adversely affects the mechanical properties of these alloys [2]. Increasing the high temperature use of near ␣-Ti alloys is being attempted by application of oxidation resistant coatings. Several coatings such as TiAlN, Al3 Ti type aluminides, silicides and MCrAlY type coatings have been reported for this purpose [3–14]. Al3 Tibased diffusion aluminide coatings formed via pack aluminizing method have been reported to greatly enhance the oxidation resistance of IMI-834 alloy up to 800 ◦ C [13,14]. Both unmodified and Pt-modified Al3 Ti coatings have been evaluated for their oxidation resistance on the above alloy and have been found to provide excellent high temperature protection to the alloy [13,14]. These coatings derive their oxidation resistance from their ability to form an adherent layer of Al2 O3 on the surface, which, once formed, retards further oxidation [14–16]. Despite their good oxidation resistance, Al3 Ti coatings are inherently brittle and develop through-thickness cracks during coating formation as well as during cyclic oxidation exposure [13,17]. Therefore, these coatings can potentially degrade the mechanical properties of the substrate alloy. Examples of brittle diffusion
∗ Corresponding author. Tel.: +91 40 24586839. E-mail address: [email protected] (C. Parlikar). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.10.021
aluminide coatings such as -NiAl degrading the mechanical properties of substrate Ni-base superalloys can be found in the reported literature [18–20]. The objective of the present study is to evaluate the effect of Al3 Ti diffusion aluminide coating on the tensile properties of a near ␣-Ti alloy at room temperature (RT) and 600 ◦ C. The effect of coating thickness on the tensile properties has also been examined in this study. 2. Experimental details An experimental near ␣ Ti-alloy having a nominal composition (in wt.%) as Ti–5.93Al–4.08Sn–3.93Zr–0.71Nb– 0.59Mo–0.37Si–0.076C, was used as the substrate material for the present study. The dissolved oxygen and nitrogen contents in the alloy were 826 and 32 ppm, respectively. The extruded alloy rods (23 mm diameter) were given the required heat treatment, i.e. solutionized at 1000 ◦ C for 2 h followed by aging at 700 ◦ C for 2 h. After both these treatments, the rods were cooled outside the furnace to room temperature in ambient air. The solutionizing and aging treatments were also carried out in air. The volume fraction of equiaxed primary ␣ in the heat treated alloy was approximately 0.3. Similar near ␣ alloys having a relatively high volume fraction of equiaxed primary ␣ phase (0.25–0.3) have been employed to achieve a good combination of creep and fatigue properties [21,22]. Cylindrical blanks of 75 mm length and 8 mm diameter were removed from the heat treated rods using electro-discharge machine (EDM) technique. Round tensile specimens having a gage length of 28 mm, parallel length of 20 mm and gage diameter of 4 mm, as per ASTM standard E8M [23], were then fabricated from the blanks in a computerized numerical-controlled (CNC) lathe
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Fig. 1. Cross-sectional microstructure of the Al3 Ti coating formed at an aluminizing temperature of 700 ◦ C. Through-thickness cracks in the coating can be seen.
machine. Some of the tensile specimens were applied with an aluminide coating by pack aluminizing method. The powder pack used for aluminization consisted of (in wt.%) 15% Al as the Al source, 2% NH4 Cl as the activator and 83% calcined Al2 O3 as the filler. The pack aluminizing process was carried out in a tube furnace. In order to obtain coatings of different thicknesses, aluminization was carried out at three different temperatures, namely 600, 650 and 700 ◦ C. The duration of aluminizing was kept constant at 5 h in all the cases. To prevent oxidation during the process, an Ar atmosphere was maintained in the tube furnace. The coated specimens, along with the uncoated ones, were tensile tested at room temperature (RT) and 600 ◦ C by using a universal testing machine (model Instron 5500R). All the tensile tests were carried out at an initial strain rate of 6.7 × 10−4 s−1 . For testing at 600 ◦ C, the tensile specimen was first fitted to the grips in the machine and then surrounded by a split-type furnace. The temperature of the furnace was then raised from RT to 600 ◦ C, which took about 1 h. During the heating-up period, the sample remained inside the furnace. After the furnace reached the above temperature, the sample was soaked at this temperature for approximately 15 min to attain temperature uniformity. Subsequently, it was tensile tested in a similar manner as mentioned earlier. At least two samples were tested under any given condition for ensuring consistency in the results. Coating microstructure was examined
Fig. 2. Composition of Al and Ti across the thickness of the coating formed at 700 ◦ C.
Fig. 3. Tensile properties of the uncoated and coated alloy: (a) at RT and (b) at 600 ◦ C.
with the help of a Quanta 400 scanning electron microscope (SEM) in as-coated condition as well as after tensile test. The phase constitution of the coating was determined by X-ray diffraction (XRD) method. A Cameca SX-100 electron probe micro-analyzer (EPMA) was utilized to analyze the composition of the coating. To examine surface cracking, the gage length portion of the failed tensile samples was observed in the above SEM. The fracture surfaces of the failed samples were also examined in the SEM.
Fig. 4. Engineering stress–strain plots for UC and C35 samples at both RT and 600 ◦ C.
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Table 1 Details of all the tensile test specimens used in the present study. Sample ID
Sample details
UC C5 C18 C35 UC5 UC18 UC35
Uncoated samples prepared from fully heat treated rods Samples with 5 m thick coating Samples with 18 m thick coating Samples with 35 m thick coating Tensile specimen prepared after machining out a 0.2 mm thick surface layer from the parallel length of a C5 sample Tensile specimen prepared after machining out a 0.2 mm thick surface layer from the parallel length of a C18 sample Tensile specimen prepared after machining out a 0.2 mm thick surface layer from the parallel length of a C35 sample
3. Results 3.1. As-coated microstructure The cross-sectional microstructure of the as-deposited coating corresponding to aluminization temperature of 700 ◦ C is shown in Fig. 1. The coating was about 35 m thick and had a single layer consisting of Al3 Ti phase, as determined from the XRD analysis [24,25]. The above phase constitution of the coating was also confirmed by EPMA chemical analysis, as shown in Fig. 2. The atom ratio of Al to Ti as 3:1 indicates the coating phase to be Al3 Ti. The coatings formed at 600 and 650 ◦ C also had the identical single-layer microstructure and consisted of Al3 Ti phase, although their thicknesses were lower at 5 and 18 m, respectively. Increase in coating thickness achieved by increasing the aluminizing temperature was consistent with the corresponding increase of Al intake by the sample during coating process [24,25]. The Al pick-up registered by the tensile specimens for the above three temperatures was 1.42, 4.63 and 8.1 mg cm−2 , respectively. Several through-thickness cracks were present in the as-deposited coatings (Fig. 1). These cracks had
formed due to the mismatch in the coefficient of thermal expansion (CTE) between the Al3 Ti coating (CTE = 15 × 10−6 ◦ C−1 [26]) and the Ti-alloy substrate (CTE = 8.9 × 10−6 ◦ C−1 [27]). Tensile stresses that were generated in the coating due to the above CTE mismatch during cooling of the coated sample after the coating process, caused the above mentioned through-thickness cracking in the brittle Al3 Ti coating [13,17,24,25]. 3.2. Tensile properties The tensile properties of the uncoated and coated alloy were determined from the engineering stress–strain plots. The properties at RT are presented in Fig. 3(a). While the uncoated tensile specimens have been referred to as UC in this study, the coated ones with coating thicknesses of 5, 18 and 35 m have been referred to as C5, C18 and C35, respectively, as indicated in Table 1. The 0.2% yield strength (YS) of the uncoated alloy was 1098 MPa and that for C5 was very similar at 1106 MPa. The YS of C18 was, however, 3% lower than that for UC alloy and the same for C35 was lower by about 12%. Thus, it is evident that, with increase in coating
Fig. 5. Slip patterns generated on the surface of the uncoated samples during the tensile test: (a) at RT (low magnification image), (b) at RT (high magnification image) and (c) at 600 ◦ C. Arrows indicate cracks in (c).
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Fig. 6. Surface morphology of the coated samples in the gage length portion: (a) C5 prior to tensile test at RT, (b) C35 after test at RT, (c) C35 after testing at 600 ◦ C. Arrows indicate cracks in (a) and (b).
thickness, the YS of the coated alloy progressively decreased. Similarly, a progressive decrease in the ultimate tensile strength (UTS) of the alloy, although marginal (<5%), was also observed as the coating thickness was increased. This is evident from the UTS values of UC, C5, C18 and C35 which were 1194, 1175, 1159 and 1141 MPa, respectively (see Fig. 3(a)). Despite causing a decrease in strength (both YS and UTS), the presence of the coating did not cause any appreciable decrease in the plastic strain to fracture of the alloy (henceforth referred to as ductility) in case of C5 and C18, as the ductility values of about 10% for these samples were similar to that of UC. However, the ductility of C35 was somewhat lower at 8% as compared to the value of 12% measured for UC (Fig. 3(a)). It has been reported that precipitation of ordered Ti3 Al (␣2 ) phase and silicides occurs in near ␣ Ti-alloys during the aging treatment [22,28,29]. The precipitated ␣2 and silicide particles and their coarsening are believed to cause some degree of increase in the tensile strength and reduction in the ductility of the alloy both at RT and high temperatures [22,30]. The present substrate alloy, in case of the coated samples, was subjected to the treatment of 5 h at 600/650/700 ◦ C during the coating formation process. This treatment was in addition to the aging treatment (2 h at 700 ◦ C) given to the alloy prior to coating formation. To examine whether the possible changes in the alloy microstructure (in terms of precipitation/coarsening of ␣2 and silicide particles [28,31]) because of the above additional exposure during coating formation has caused any change in the tensile properties of the alloy, the coating was completely removed from a set of coated tensile samples by machining out a 0.2 mm thick layer from the surface over the entire parallel length. These samples have been referred to as UC5, UC18 and
UC35, signifying the initial thickness of the coating that was present on these samples before its removal (Table 1). These samples were then tested at RT in a similar manner as described earlier for C5, C18 and C35. The tensile properties of these samples are also shown in Fig. 3(a). It is clear that strength values (both YS and UTS) of UC5, UC18 and UC35 were very similar to each other (within ±5%) and to those of UC. The ductility values of UC5, UC18, UC35 were also very similar to that of UC and lied in the range 10–13%. This observation clearly indicates that the high temperature exposure during the coating treatment, i.e. 5 h at 600/650/700 ◦ C, did not cause any appreciable change in the tensile properties of the substrate alloy. Therefore, the minor degradation in the tensile properties of the alloy observed in the coated condition, i.e. in samples C5, C18 and C35, as mentioned previously, was primarily caused by the presence of the coating. Similarly, the somewhat lowering of ductility in case of C35, as mentioned earlier, was also caused by the coating, although such an effect was not prominent in case of C5 and C18. The tensile properties of the present alloy at 600 ◦ C are shown in Fig. 3(b). The YS and the UTS for the uncoated (UC) alloy at the above temperature were 597 and 721 MPa, respectively. Similar to that observed at RT, a reduction in alloy strength (YS and UTS) was also noticed at 600 ◦ C. The trend of progressive decrease in YS with increase in coating thickness was also observed at 600 ◦ C, as evident from the YS values of 580, 567 and 523 MPa measured for C5, C18 and C35, respectively (Fig. 3(b)). A similar trend was also maintained in the UTS values of the coated alloy (Fig. 3(b)), with UTS of C35 at 590 MPa, being nearly 20% lower than that of UC. It is evident from Fig. 3(b) that the ductility value of 11.5% measured for UC at 600 ◦ C was very similar to that measured at
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RT (Fig. 3(a)). However, the coated alloy showed a much higher ductility at 600 ◦ C than at RT. The ductility of the coated alloy at 600 ◦ C lied in the range 14.5–20.5%, which was significantly higher than the corresponding value of 11.5% for UC. Further, among the coated samples, a decrease in the ductility value from about 20% to 14.5% was observed as the coating thickness increased from 5 m to 35 m. As was done at RT, the coating-removed samples UC5, UC18 and UC35 samples were also tested at 600 ◦ C and the tensile properties of these samples are shown in Fig. 3(b). As clearly evident, YS, UTS and ductility of UC5, UC18 and UC35 were very similar to the corresponding values of UC. This again suggests that the degradation in strength and the enhancement in ductility observed for the coated alloy at 600 ◦ C were primarily caused by the presence of the coating. The typical tensile curves for uncoated and coated samples corresponding to RT and 600 ◦ C are shown in Fig. 4. The curves for the coated sample in the above figure are for C35. As typical of near ␣ Ti-alloys [32], the RT tensile curve for UC did not show any appreciable sign of post-UTS necking till failure. The behavior of the RT tensile plot for C35 sample was also very similar to that of UC. Both uncoated and the coated samples exhibited enhanced post-UTS necking when tested at 600 ◦ C, as evident from decreasing stress with increasing strain beyond UTS in Fig. 4. The tensile behavior of the present near ␣ Ti alloy in uncoated and coated conditions can be summarized as follows. Application of Al3 Ti coating on the alloy led to a marginal (up to 12%) reduction in strength (both YS and UTS) at RT, with the extent of reduction increasing with the increase in coating thickness. The presence of the coating did not cause any appreciable reduction in alloy ductility at RT for C5 and C18, which remained in the range of 10–12%. Only in the case of C35, the ductility reduced to 8% as compared to the value of 12% for the uncoated alloy. The variation of strength of the coated alloy at 600 ◦ C was similar to that observed at RT. However, unlike at RT, the presence of coating led to an enhancement in the ductility of the alloy at 600 ◦ C, although a trend of decrease in ductility with increase in coating thickness was observed at this temperature. 3.3. Microstructural and fractographic observations in tensile samples 3.3.1. Surface features over gage length Typical surface features on the gage length portion of the failed UC samples are shown in Fig. 5(a). Numerous near-parallel lines were found criss-crossing the surface of the tested samples. These marks were not pre-existing on the as-machined samples and were generated during the tensile test. A magnified view of the surface, as shown in Fig. 5(b), suggests the above marks to be sets of parallel slip lines distributed over the sample surface. The angles between the intersecting sets of parallel slip lines were measured to be 53 −
−
and 55◦ . It has been reported that usually {0 0 0 1}, {1 0 1 0}, {1 0 1 1} −
and {1 1 2 2} slip planes operate during deformation in ˛/ˇ Ti alloys [33]. The interplanar angle () between these slip planes can be calculated by using the following equation (Eq. (1)) [34]: 2
hd + ke + (1/2)(he + kd) + (3/4)lg(a/c)
cos =
2 1/2
{h2 + k2 + hk + (3/4)l2 (a/c) }
2 1/2
(1)
× {d2 + e2 + de + (3/4)g 2 (a/c) }
The indices (hkil) and (defg) in the above equation denote the Miller indices for two planes between which the angle is being calculated. The measured angles, as mentioned above, corresponded to comparison of the above measured values of 53 and 55◦ with various interplanar angles indicated that the participating slip planes in the −
deformation of UC sample during tensile test were {0 0 0 1}, {0 1 1 1} −
Fig. 7. Longitudinal sections of the tensile specimens tested at RT: (a) UC, (b) C5 and (c) C35. Magnified view showing voids is presented in the inset in (a).
and {1 0 1 0} type. Same slip line pattern was also observed on the UC sample tested at 600 ◦ C, as evident from Fig. 5(c). Additionally,
this sample showed extensive surface cracking with the cracks oriented transverse (perpendicular) to the loading direction, as shown in the above figure. Cracking along slip lines was also observed at several locations. A few fine cracks could be located on the surface of as-coated specimens (C5, C18 and C35), as typically shown in Fig. 6(a) for C5 sample. These cracks correspond to the through-thickness cracks generated in the coating under the influence of CTE mismatch, as
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Fig. 8. Longitudinal sections of the tensile specimens tested at 600 ◦ C: (a) UC, (b) C5 and (c) C35.
mentioned earlier (Fig. 1). During tensile test at RT, many additional cracks were generated in the coating transverse to the loading direction, as typically shown in Fig. 6(b) for C35. The cracks generated during RT testing largely remained fine in all the three coated samples. No slip traces, as seen in UC sample, were observed on the surface of the failed coated samples. The coated samples tested at 600 ◦ C also revealed large number of transverse cracks on the surface, as typically shown in Fig. 6(c) for failed C35 sample. These cracks were also much wider than those generated by testing at RT (Fig. 6(b)). 3.3.2. Observations on longitudinal-sections The longitudinal sections of the failed UC, C5 and C35 samples, tested at RT, are presented in Fig. 7(a)–(c). It is evident from Fig. 7(a) that no cracking developed at the surface, either close to fracture end or away from it, prior to the failure of UC sample. This is unlike the case in C5, C18 and C35 where many through-thickness cracks of the coating could be found penetrating into the substrate over the entire gage length, as typically shown in Fig. 7(b) and (c) for C5 and C35, respectively. These through-thickness cracks could also be seen on the sample surface, as mentioned previously (Fig. 7(b)). In case of the uncoated sample, several voids were found to be present in the substrate close to the fracture end (Fig. 7(a)). A closer examination revealed that the voids had formed by de-cohesion of the ␣ lath boundaries within the transformed  colonies. Fig. 8 presents the longitudinal section of a failed UC sample along with those of C5 and C35 samples tested at 600 ◦ C. Many surface cracks, penetrating into the substrate could be observed, throughout the longitudinal section, in case of UC (Fig. 8(a)). These cracks corresponded to the transverse cracks that were seen on the surface, as mentioned earlier (Fig. 5(c)). The cracks seen in the longitudinal section of UC, largely remained at the surface and did not
penetrate to any appreciable depth into the sample, except close to the fracture end. In case of the three coated samples, throughthickness cracks were present in the coating all along the gage length and some coating cracks were found to have penetrated into the substrate, especially in C35 (Fig. 8(c)). Apart from the cracks at the surface, voiding in the substrate close to the fracture end was observed in both uncoated and coated samples that were tested at 600 ◦ C. In order to understand the effect of cracks in the coating on the strength of the alloy, especially UTS, a few tests were carried out at RT and 600 ◦ C wherein the loading was stopped very close to UTS. Subsequently, the samples were unloaded and their longitudinal sections metallographically polished and inspected in SEM. A large number of through-thickness cracks along the gage length could be found in all the coatings at both the temperatures, similar to that mentioned earlier for the failed specimens (Fig. 9(a)). Further, the penetration of the cracks into the substrate was also observed in all the three cases at both the temperatures. However, at 600 ◦ C, the number of penetrating cracks and their penetration depths into the substrate were much higher for C35 than for either C5 or C18, as evident in Fig. 9(b). In fact, most of the cracks in the latter two cases remained restricted to within the coating (Fig. 9(a)). 3.3.3. Fractography Fig. 10(a) presents the fracture surface of a failed UC sample tested at RT. The presence of quasi-cleavage features over the entire fracture surface was observed in this sample. The fracture surface of all the coated specimens tested at RT, including C35, also exhibited similar quasi-cleavage features, as typically shown in Fig. 10(b). The separation along the interfaces of ␣ laths in the transformed  colonies could be clearly seen on the cleavage planes, as shown in Fig. 10. Unlike at RT, the failure at 600 ◦ C
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Fig. 9. Longitudinal sections of the tensile specimens after interrupted tensile testing at 600 ◦ C: (a) C5 and (b) C35.
occurred in a ductile manner in all the four cases (UC and the three coated samples). This was evident from the dimpled appearance of the fracture surfaces for these samples, as typically shown in Fig. 11(a) for UC and Fig. 11(b) for C35. The presence of several voids in the substrate close to the fracture end in all the failed samples, as mentioned earlier, supports the above ductile failure mechanism involving the nucleation of voids and their subsequent coalescence causing fracture. 4. Discussion As mentioned previously, the substrate alloy underwent an additional exposure for 5 h at the aluminizing temperature (600/650/700 ◦ C), which was close or equal to its aging temperature. It has been reported that, in near ␣ alloys, some degree of precipitation and subsequent coarsening of Ti-silicides and ␣2 (Ti3 Al) phase occur during the aging treatment as well as subsequent use at high temperatures [22,28,29]. Zhang et al., in case of two near ␣ alloys, have shown that such precipitation and/or coarsening during the aging treatment did not cause any appreciable change in alloy strength and ductility, irrespective of the duration of aging (which was varied in the range 2–15 h at aging temperatures of 700 and 760 ◦ C) [22]. Only when long term thermal exposure at 600 ◦ C for 100 h was given, an increase in alloy strength and a decrease in ductility was observed, which was ascribed to additional precipitation/coarsening of silicides and ␣2 particles. Woodfield et al. have also made similar observations in case of near ␣ alloy Ti 5331S [30]. Based on these reported studies, it is clear that
Fig. 10. Fracture surfaces of the failed tensile samples tested at RT: (a) UC and (b) C35. Cleavage separation along the ␣ lath boundaries within transformed  colonies can be seen from the inset in Fig. 9(b).
the effect of aging temperature and its duration on the strength and ductility of near ␣ alloys is at best limited. Therefore, it can be assumed that the additional high temperature exposure of the present substrate alloy during the aluminizing treatment would not cause any appreciable change in its tensile properties. This aspect is also clearly evident from the fact that the strength and ductility values of UC, UC5, UC18 and UC35 were very similar at both RT and 600 ◦ C (Fig. 3). Being a diffusion coating, the present Al3 Ti coating formed an integral part of the substrate. As mentioned earlier, the brittle coating, irrespective of its thickness, had several through-thickness cracks in as-formed condition (Fig. 1). Since Al3 Ti phase remains brittle up to 800 ◦ C [35,36], additional cracks were generated in the coating during very early stages of tensile loading at both the test temperatures (Figs. 7 and 8). Because of the presence of the through-thickness cracks, the coating in the coating-substrate composite structure, i.e. in the coated samples, can be considered as non-load bearing during the tensile test. However, the coating cracks would cause stress concentration at the crack tips, i.e. in the coating/substrate interface. Because of stress concentrations at the crack tips, the substrate would locally yield at a stress lower than
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Fig. 11. Fracture surfaces of the failed tensile samples tested at 600 ◦ C: (a) UC and (b) C35.
the YS of the uncoated alloy. Thus, the YS of the alloy would decrease in coated condition, as was observed in the present study (Fig. 3). The extent of YS drop in the coated samples would depend on the magnitude of the stress concentration at the crack tips, which is function of the crack length (or coating thickness) and crack tip radius. Assuming the crack tip radius to be equal in case of all the three coated samples in as-formed condition, the stress concentration at the crack tips would have a direct dependence on the crack length [37]. Thus, the decreasing trend in the YS observed in the coated samples with increase in coating thickness was due to the corresponding increase in the stress concentration ahead of the through-thickness crack tips present in the coatings. There are two major factors that are likely to cause a reduction in the UTS values of the coated samples: (1) the non-load bearing nature of the coating and (2) the penetration of the coating cracks into the substrate leading to further lowering of the load bearing cross-section of the sample. Thus, it can be concluded that the reduction in the UTS values of the coated samples observed in this study (with respect to the UTS of UC) has been a combined contribution of the above two factors. To evaluate the contribution of the non-load bearing coating, the coating thickness (t) can
be subtracted from the initial diameter (D) of the coated sample and the UTS can be revised upward based on the reduced diameter, i.e. D − 2 × t. The revised UTS values for C5, C18 and C35 at RT are 1181, 1180 and 1182 MPa, respectively. These values are very similar (within ±2%) to 1194 MPa measured for the UC sample (see Fig. 3(a)). Thus, it is clear that the presence of the non-load bearing coating was primarily responsible for causing a reduction in the UTS of the coated samples at RT, with the degree of reduction directly dependent on the coating thickness. The contribution of the second factor, i.e. the lowering of the load bearing cross-section of the substrate because of the penetration of coating cracks, to the above reduction in UTS was minimal. The UTS values for the coated samples at 600 ◦ C can also be revised upward by subtracting the coating thickness from the sample diameter, as mentioned above. The revised UTS values at 600 ◦ C for C5, C18 and C35 are 651, 649 and 611 MPa, respectively. These revised values, which include only the contribution of the nonload bearing nature of the coating, are 10–15% lower than the UTS value of 721 MPa measured for UC sample (Fig. 3(b)). This difference between the revised values and the UTS of uncoated sample has actually been contributed by the lowering of the load bearing cross-section of the substrate because of the penetration of coating cracks. The decreasing trend in the UTS of coated samples at 600 ◦ C with increasing coating thickness was in agreement with the previously mentioned observation that the crack penetration depth into the substrate increased with coating thickness during the tensile test (Fig. 8). The effect of through-thickness cracks in the coating and their penetration into the substrate, as mentioned above, did not cause any significant reduction in the ductility in C5 and C18 specimens at RT (Fig. 3(a)). However, the somewhat lower ductility of C35 as compared to UC was possibly because of the much higher crack length and the corresponding higher stress concentration ahead of the cracks in the coating. From Fig. 10(a), it appears that the decrease in ductility in case of C35 was too small to be reflected in terms of any change in fracture surface features. The uncoated alloy was expected to show a comparatively higher ductility at 600 ◦ C than at RT. However, the ductility of UC at both temperatures was found to be virtually the same (Fig. 3). As stated previously, for testing at 600 ◦ C, the tensile specimen remained inside the furnace for about 1.25 h during which the furnace attained the above temperature and homogenization of temperature occurred inside the sample. It is believed that, during this period, the surface of the sample underwent some degree of oxidation-related damage, possibly in terms of the formation of a thin ␣ case. Although the presence of an ␣ case could not be confirmed, the several surface cracks observed in the failed sample, as seen in Fig. 5(c) and Fig. 8(a), supports the possibility of the above surface damage to the UC alloy. Thus, even though the alloy inherently has a higher ductility at 600 ◦ C, it was most likely negated by the loss in ductility caused by the above surface damage [2,21]. Consequently, no significant increase in ductility was registered for UC at 600 ◦ C as compared that at RT. In case of the coated specimens, the presence of the protective Al3 Ti coating prevented the substrate alloy to undergo any surface damage during the heating-up period prior to tensile testing. As a result, much higher ductility was obtained in these samples at 600 ◦ C as compared to that in UC. Such increased ductility values of the coated samples are expected to be close to the actual ductility of the substrate alloy at 600 ◦ C, in the absence of surface damage caused by oxidation. It is interesting to note that such increased ductility for the coated alloy was observed at 600 ◦ C despite the presence of numerous through-thickness cracks in all the three coatings (Fig. 6(c)). Increase in the coating thickness and, hence, the length of the through-thickness cracks (Fig. 3(b)) did cause some drop in the ductility of the alloy, as evident from Fig. 3(b). However, the enhancement in ductility resulting from the surface protection
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provided by the coating was much larger to negate the effect of the coating cracks. As a result, an appreciable increase in the overall ductility was registered by the coated alloy at 600 ◦ C for all three coating thicknesses (Fig. 3(b)). The beneficial effect of the coating, as described above, has also been observed by Leyens et al. [38] in their creep study (at 600 ◦ C) on uncoated and aluminide coated near ␣ alloy Timetal 1100. From the present study, it also appears that the cracking in the coating does not significantly affect the ductility of the alloy even at RT. The results of the present study suggests that the application of the diffusion aluminide (Al3 Ti) coating does not cause any significant loss in ductility of the experimental near ␣ Ti-alloy both at RT and 600 ◦ C. However, it leads to a marginal (10–18%) decrease in the alloy strength. From the tensile results of this investigation and the reported excellent oxidation resistance of Al3 Ti coating [13,14], it can be concluded that the above coating shows good promise for providing high temperature oxidation protection to near ␣ Ti-alloys, without causing appreciable degradation in their tensile properties. 5. Conclusions The tensile properties of a near ␣ Ti-alloy have been evaluated in uncoated as well as diffusion aluminide (Al3 Ti) coated conditions at RT and 600 ◦ C. The effect of coating thickness (5, 18 and 35 m) on the tensile properties of the alloy has also been examined. The presence of the coating was found to marginally reduce (by up to 12%) the YS and UTS of the alloy at RT, with the extent of reduction increasing with the increase of coating thickness. Such a trend of decrease in alloy strength with increase in coating thickness was also observed at 600 ◦ C. The effect of coating thickness on RT ductility of the alloy was marginal. Only in the case of highest coating thickness (35 m), the RT ductility reduced to 8% from the value of 12% for the uncoated alloy. However, an enhancement in ductility of the alloy was observed in presence of the coating at 600 ◦ C for all the three coating thicknesses. The drop in YS observed in the coated alloy was explained in terms of the stress concentration caused by the through-thickness cracks present in the coating. The drop in UTS, on the other hand, could be ascribed to the combined effect of non-load bearing nature of the coating and the loss in load bearing cross-section because of the penetration of coating cracks into the substrate. Acknowledgements The authors acknowledge the assistance provided by MBG, SFAG and EMG groups of DMRL. They are thankful to Director, DMRL, for his permission to publish the present work. This research work has been funded by the Defence Research and Development Organization, DRDO.
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