Firstly, the macrostructure on the longitudinal section, stretching from the thick section (19 mm) to thin section (4 mm), of the cylindrical casting has been analyzed, as shown in Fig. 2. The outlined prior β grain morphology shows that, for both thick and thin section, the edge part tended to have a columnar prior β grain morphology and the centre part tended to have an equiaxed morphology. The grain morphology is controlled primarily by the local thermal conditions^[2]. During the solidification of the liquid metal in the thin walled cavities, the direction of heat flow was from the centre to the edge. On the mould-metal interface, the β grain firstly nucleated and then grew along the negative direction of the heat flow. In the centre part, a homogeneous nucleation and growth of β grains occurred. Besides the β grain morphology, it was also found that the β grains in thick section are overall larger than those in thin section, and the maximum prior β grain size is about 3.5 mm in diameter. Meanwhile, the fine-scale microstructure varies with both the location and section size, as shown in Fig. 3. Compared with the thin section, thick section has larger α /β colonies. And the centre part has larger α /β colonies than the edge part. The fine-scale microstructural morphology is controlled primarily by the post-solidification cooling rate.

Fig. 2.

	Figure Option ViewDownloadNew Window
	Fig. 2. Micrographs showing the macrostructure on the longitudinal section, stretching from the thick section (19 mm) to thin section (4 mm), of the cylindrical casting with a mould pre-heating temperature of 873 K and after HIPping. OM image (left) and the outlined prior β grain morphology (right).

Fig. 3.

	Figure Option ViewDownloadNew Window
	Fig. 3. OM images showing the microstructures at different locations of the macrograph in Fig. 2: (a) edge of thin section, (b) centre of thin section, (c) edge of thick section, (d) centre of thick section.

In addition to the local microstructural inhomogeneity, mould pre-heating temperatures and HIPping process also affect the microstructural characteristics, as shown in Fig. 4 (Note: the casting with mould pre-heating temperature of 873 K was not cut in half before HIPping process). For the pre-heating temperature of 673 K, the casting consists of a mixture of Widmanstä tten and colony microstructure. With increasing pre-heating temperature, the casting is consisted primarily of colony α microstructure. By comparing these three casting microstructures, it is shown that the structural unit size increases with increasing the mould pre-heating temperature. The lower mould pre-heating temperature brought about relatively higher cooling rate, thus giving rise to a much finer microstructure. The major effects of HIPping process on the casting microstructure are the removal of porosities and promotion of grain growth^[5]. After the HIPping process, as shown in Fig. 4, both α /β colonies and α laths became coarsened due to the combined effect of the high pressure and temperature. The condition at lowest preheating temperature was the most obvious. The mould pre-heating temperature also affects the casting defects, especially for the porosities, as shown in Fig. 5. With increasing the mould pre-heating temperature, the number of porosities decreases dramatically. Under the casting condition with low mould pre-heating temperature, the solidifying time of the casting is short. Entrapped gas could not be expelled before the thin walled casting is solidified, so there are more casting pores in the casting^[12].

Fig. 4.

	Figure Option ViewDownloadNew Window
	Fig. 4. OM images showing the effects of mould pre-heating temperatures and HIPping process on the casting microstructures: (a) 673 K, (b) 673 K + HIP, (c) 773 K, (d) 773 K + HIP, (e) 873 K + HIP.

Fig. 5.

	Figure Option ViewDownloadNew Window
	Fig. 5. X-ray detection showing the variation of casting defects with the mould pre-heating temperatures: (a) 673 K, (b) 773 K. The dark spots represent the casting porosities, as indicated by white arrows. The casting porosities from 773 K have been presented with (c) OM image from the cross section and (d) SEM image from the tensile fracture.

Factors influencing the casting microstructure mainly include the chemical composition, pouring temperature and cooling rate during and after solidification. The chemical composition and the pouring temperature are the same for all the castings in the present study. The cooling rate of the investment castings depends on the heat transfer between the mould material and casting metal, which is related to the heat transfer coefficient and heat capacity of mould materials, possible reaction with mould materials, mould pre-heating temperature and wall thickness of the casting. In this study, obviously, the last two factors determine the microstructural characteristics. In thin section, the heat could be transferred to the mould more quickly than that in thick section. Thus the cooling rate in thin section is higher than that in thick section. The grain size obtained during solidification is determined by the competition between the nucleation and growth rates. With increasing cooling rate, the nucleation ratio increases. But the diffusion rate decreases, so the new phases have difficulty in growing. The combined effect of the increase in nucleation ratio and decrease in growth rate results in smaller grain size^[3].

The tensile properties of the Ti-6Al-4V castings are listed in Table 1. By comparing these five different processing conditions, it can be found that, for the as-cast condition, both the tensile strength and elongation of the casting with the mould pre-heating temperature of 773 K are higher than those at 673 K. However, for the condition after HIPping, an inverse situation emerges. With increasing mould pre-heating temperatures, both the tensile strength and elongation of post-HIPped castings decrease.

Table 1

Table 1

Table 1 Tensile properties of the thin sections from the cylindrical castings produced under different processing conditions (average of three test specimens) Processing conditions YS (MPa) UTS (MPa) Elongation (%) Mean S D Mean S D Mean S D 673 K 742.3 23.7 827.0 31.1 3.9 0.5 673 K + HIP 811.5 21.4 890.7 36.4 6.1 0.4 773 K 782.7 18.6 865.3 6.1 4.8 0.2 773 K + HIP 782.8 15.5 840.6 31.3 4.9 0.5 873 K + HIP 767.2 26.2 832.3 11.0 4.7 0.2

Table 1 Tensile properties of the thin sections from the cylindrical castings produced under different processing conditions (average of three test specimens)

To clarify the variation of tensile properties as stated above, the fracture surfaces of the tensile specimens were carefully examined and compared with the aid of SEM, taking the casting of 673 K before and after HIPping as the example, which is shown in Fig. 6. Obviously, there are two distinct fracture surfaces. For the as-cast sample, the failure is typically dominated by the intergranular cleavage fracture where most of the decohesion appear at prior β grain boundary, and cast pores can be easily found on the fracture surface, as indicated in Fig. 6(c). But for the post-HIPped sample, it is a typical ductile fracture and associated with more dimples and mostly inter-colony decohesion. On the whole, the fracture surface of post-HIPped sample is relatively flatter than that of the as-cast. The variation of fracture features before and after HIPping process directly explains the increase of tensile strength and elongation after HIPping at 673 K. Specifically, the HIPping process remarkably eliminates the cast pores and improves the cohesive strength of β grain boundary. With increasing the mould pre-heating temperature from 673 to 773 K, both the strength and ductility of the as-cast sample are increased. This is mainly because the higher pre-heating mould temperature would decrease the solidification rate and thus reduce the presence of cast pores, as shown in Fig. 5. Although the higher mould temperature would produce the relatively coarser microstructure and decrease the strength, it played a secondary role while comparing with the positive effect of the elimination of cast pores. On the improvement of the tensile performance, HIPping process had a more obvious effect for the lower pre-heating temperature case, because the number and size of micro-pores would become much smaller in the castings with increasing mould pre-heating temperature.

Fig. 6.

	Figure Option ViewDownloadNew Window
	Fig. 6. Comparison of the SEM tensile fractographs of the casting with mould pre-heating temperature of 673 K before and after the HIPping processing: (a, c) as-cast, (b, d) post-HIPped.

For the post-HIPped samples, the strength and ductility have a decreasing tendency with the increase of mould pre-heating temperatures due to the microstructure coarsening. For the lower pre-heating temperatures, the casting microstructure is much finer, as shown in Fig. 4. The finer microstructural size of Ti-6Al-4V alloy castings increases both the grain and phase boundaries, and then results in the higher strength and ductility, which can be explained by the Hall-Petch mechanism. This viewpoint is also supported by the observation that the strength tends to increase as the colony size decreases^{[8, 13 , 14]}.

The directionality of heat transfer and non-uniformity of cooling during the casting process of thin-wall products determine the inhomogeneous microstructures and properties in the whole cross section. Fig. 7 shows the variation of the Vickers micro-hardness from surface to inner under five processing conditions. In most cases, the micro-hardness reduces rapidly from the surface to a depth of 240 µ m and then reduces slowly to a depth of 450 µ m. When the depth is more than 450 µ m, the micro-hardness tended to be a constant. Thus, the thickness of the reaction layers is about 450 µ m for the mould in this study. The variation tendency of the micro-hardness and the thickness of the reaction layer are almost the same for the investigated five processing conditions. It can also be seen that, on the whole, the micro-hardness decreases with the increase of mould pre-heating temperature. The casting microstructure obtained at the mould pre-heating temperature of 673 K has the highest Vickers micro-hardness.

Fig. 7.

	Figure Option ViewDownloadNew Window
	Fig. 7. Micro-hardness profiles near the surfaces of the thin sections of the cylindrical castings produced by different processing conditions.

The reason for the variation of micro-hardness stated above can be attributed to two aspects. First, the cooling rate during the solidification presents a descending change from the surface to the inner of castings. So, the microstructure has a corresponding variation from fine to coarse. The Vickers hardness decreases with the increase of microstructure size. The smaller microstructure produces higher grain and phase boundary density and thus higher Vickers hardness. The higher mould pre-heating temperature also corresponds to lower cooling rate and coarser casting microstructure. Second, the surface hardening layer is caused by the oxygen contamination^[3].

To satisfy the requirement of the application in aerospace field, the damage-tolerance properties of the Ti-6Al-4V cylinder casting, including fracture toughness and FCG, have been evaluated in this and in the next section, respectively. Because the thicknesses of the extracted CT samples, which is restricted by the thickness of the casting wall, cannot satisfy the fracture toughness test standard for metals^[15], the values of mode I fracture toughness, K_q, of the casting process under different processing conditions are shown in Fig. 8. It should be noted that the K_q values calculated on the basis of a plane strain concept are usually higher than K_{Ι C} values. For the engineering application, however, the information given by the K_q values is still sufficient for the property evaluation and comparison, because they represent the force to trigger unstable crack propagation for the cylinder castings investigated here. As shown in Fig. 8, the fracture toughness of both the as-cast and post-HIPped Ti-6Al-4V castings increases with increasing mould pre-heating temperature. And the post-HIPped has higher K_q values than the as-cast. The most significant improvement in fracture toughness has been found for the casting at 673 K.

Fig. 8.

	Figure Option ViewDownloadNew Window
	Fig. 8. Fracture toughness (Kq) as a function of the processing conditions (average of three test specimens).

Fig. 9 shows the fracture surfaces of the as-cast and post-HIPped samples of 673 K after the fracture toughness tests. The fracture surface images indicate that the as-cast sample mainly fractured along prior β grain boundaries, which corresponds to the lower-energy crack paths and thus led to the lower fracture toughness. Grain boundaries, as the last solidification regions during the casting process, are usually the weakest positions in titanium castings and have micropores, as shown in Fig. 9(a). Micropores in the boundaries of prior β grains are apt to cause stress concentration and develop into a crack at the micropore tip. Once the cracks nucleated from micropores in the alloy, they will propagate along the weak boundaries under the imposed stress, which aggravate the failure of the material. The fracture mode transforms from the mixed intergranular and transgranular fracture to predominately transgranular fracture after the HIPping process. The transgranular fracture with large and elongated dimples, as shown in Fig. 9(b), is mainly inter-colony decohesion. Such a situation is similar to the tensile fractures, as shown in Fig. 6. Therefore, it can be concluded that macro- and micro-porosities are removed and grain boundary cohesion is enhanced by HIPping process^[16], and then the fracture toughness of the casting microstructure will be remarkably improved.

Fig. 9.

	Figure Option ViewDownloadNew Window
	Fig. 9. Comparison of the SEM fractographs of the casting with mould pre-heating temperature of 673 K after fracture toughness tests: (a) as-cast, (b) post-HIPped.

Another variation is that the K_q values increases with increasing mould pre-heating temperatures. As we stated in Section 3.1, both the high mould temperatures and HIPping process will make the microstructure coarse. The coarser microstructures usually lead the more tortuous crack paths, which will consume more energy and increase the fracture toughness of the alloys.

FCG rates (da/dN) as a function of applied stress intensity range (Δ K) for the cylindrical Ti-6Al-4V castings produced under different processing conditions are plotted in Fig. 10(a). The tested samples were extracted from the thin section and with a thickness of 3.0 mm. It shows that with increasing the mould pre-heating temperatures, the FCG rates decrease. And after the HIPping process, FCG rates also slightly decrease.

Fig. 10.

	Figure Option ViewDownloadNew Window
	Fig. 10. FCG rates (da/dN) as a function of applied stress intensity range (Δ K), tested under R = 0.1, for the cylindrical Ti-6Al-4V castings: (a) comparison of different processing conditions, (b) effect of sample location and thickness, which are conducted on the 873 K + HIP casting.

As the solidification process and microstructure would be affected by the thickness of the cyclinder wall, some samples with different thickness from the thin and thick sections were employed to compare their FCG behaviour, as shown in Fig. 10(b). It shows that FCG data of 3.0 mm are more scattered than that of 7.5 mm for the thick section. The possible reason is that the sample from the thick section with a thickness of 3.0 mm contains few to less to least colonies in the sample thickness direction. While the crack propagate ahead, the encountered colonies have different crystallographic orientations (either favourably or unfavourably orientated for the crack growth), which will result in the considerable variation in FCG rates and then present a scattered distribution in the FCG rates^[17]. Another feature is that the sample from the thin section and with a thickness of 3.0 mm has the similar FCG resistance with the sample from the thick section with a thickness of 7.5 mm. However, the sample from the thick section with a thickness of 3.0 mm has intersected with the above two FCG rate curves and can be divided into two stages, as indicated by the letter “ A” in Fig. 10(b). Before the point “ A” , the FCG rate is lower than that of the other two samples, and the opposite case appears after point “ A” . The reason for this situation is a combined effect of structural unit size, mainly the α /β colony, and sample size. The sample from the thick section has relatively lager α /β colonies. While the sample is 3.0 mm in thickness, the relatively large and few α /β colonies will result in the FCG path having higher levels of crack deflection and roughness-induced crack closure, and thus lower FCG rates. But for the sample with 7.5 mm, more α /β colonies existing in the thickness direction means more chances to change the propagation direction and finally exhibit relatively flatter FCG path on the whole. For the later stage of the investigated FCG regime of the sample with 3.0 mm, quasi-cleavage fracture of α /β colonies simultaneously occurs on the nearly whole thickness and almost constitutes one facet in thickness direction, as shown in Fig. 11, and results in the fast propagation.

Fig. 11.

	Figure Option ViewDownloadNew Window
	Fig. 11. SEM image showing the FCG path of the sample with a thickness of 3.0 mm from the thick section of the 873 K + HIP casting. The regions, as indicated with “ A” and “ B” , correspond to the labels in Fig. 10(b). FCG direction is from left to right.

Fig. 12 shows the longitudinal section of the as-cast and post-HIPped samples after FCG tests along the fatigue crack path. It indicates that the post-HIPped sample has more crack branching, micro-cracking and deflection, and its crack path is more tortuous than that of the as-cast. In this study, the crack propagate forward mainly along a path perpendicular to the lamella long axis direction. And the quasi-cleavage of α /β colony along intense planar slip bands is the predominant failure mode^[18]. The quasi-cleavage cracking observed on the longitudinal section can be further illustrated by the planar facets on the fracture surfaces, as shown in Fig. 13. Usually an α /β colony corresponds to a planar facet. The as-cast sample has much smaller colonies, which bring about the smaller facets. And the smaller facets on the different planes connect with each other, which constitute a relatively flatter fracture surface than the post-HIPped sample.

Fig. 12.

	Figure Option ViewDownloadNew Window
	Fig. 12. OM images showing the FCG path profiles of Ti-6Al-4V castings with different processing conditions: (a) 673 K, (b) 873 K + HIP.

Fig. 13.

	Figure Option ViewDownloadNew Window
	Fig. 13. SEM fractographs of the FCG samples from the casting with a mould pre-heating temperature of 773 K: (a) as-cast, (b) post-HIPped.

The planar facets on the fatigue fracture surface are related to the planar slip behaviour in α lamellae, as shown in Fig. 14. It shows that planar slip bands include those both parallel and perpendicular to the long axis direction of α lamellae. In titanium alloys, the priority order to start the slip systems is < a > prismatic, < a > basal and < c + a > pyramidal slip, which is determined by the corresponding critically resolved shear stress (CRSS) ^{[19 , 20]}, and [0001] c-axis of α phase is quasi-parallel to the long axis direction of β platelets ^{[18, 21 , 22]}. It can be seen that, in this study, the fatigue dislocation activate along ( ) prismatic slip and (0001) basal slip, respectively, as shown in Fig. 14(a, b). The high anisotropy of the principal slip systems in the hexagonal crystal determined the high tortuosity of the crack path. The ( ) prismatic slip dislocations travel long distance to the end of α colonies and pile-up at the colony boundaries. The typical configuration of (0001) basal slip dislocations is found to be a pile-up at the α /β interface, thus imposing an additional stress on the interface. Although the occurring of planar slip is limited in α lamellae, α -Ti and β -Ti platelets assume to have perfect Burgers' relationships between ^{[21 , 23]}. The breakdown of the very first barrier enables doubling the pile-up length and finally the shear band propagates to the end of the colony^[24].

Fig. 14.

	Figure Option ViewDownloadNew Window
	Fig. 14. Bright-field TEM micrographs showing planar slip bands with respect to the α /β colony long axis direction: (a) parallel, (b) transverse.

Taking one of the casting microstructures as an example, the EBSD technique was further utilized to characterize the relationship between the crystallographic orientation of α /β colonies and planar facets on the fracture surfaces, as shown in Fig. 15. Colonies 1, 2, 3, 6 and 7 are cracked on basal plane, and colonies 6 and 8 are cracked on prismatic plane. In colony 6, the fatigue crack propagation path has transformed from along the basal slip to along the prismatic slip. Although the probability of the occurrence of prismatic slip is relatively larger than that of basal slip, and grains orientated for prismatic slip have also been reported to act as crack initiation sites^{[25] and [26]}, the ultimately cracked planes for the current case mostly appeared on the basal plane, which coincide with the previous research^{[25, 27, 28]}.

Fig. 15.

	Figure Option ViewDownloadNew Window
	Fig. 15. EBSD mapping showing the crystallographic orientation of α colonies near the FCG path. 3D crystal viewers have been presented for the corresponding labelled α colonies.

However, there are still some colonies that did not crack along favourable slip orientation, such as colony 5. Bantounas et al.^[27] explained this kind of phenomenon by assuming that all grains had active slip systems prior to cracking. The unfavourable orientated crack planes are believed to be influenced by prior cracking in neighbouring grains and are a consequence of accommodating crack propagation. For the cracks forming along favourable slip planes, once the crack propagated into a differently orientated neighbouring colony, the crack propagating direction would change into the slip band in the neighbouring colony. If the misorientation between two neighbouring colonies was large, there would be a larger deflection in the FCG path.

In conclusion, the larger α /β colony sizes and misorientation between two neighbouring colonies will cause larger tortuosity in the FCG path, which increase the crack closure degree and reduce the effective Δ K value, then decrease the crack growth rate da/dN. Thus, with increasing mould pre-heating temperatures from 673 to 873 K and after the HIPping process, the casting microstructure will become coarser and the crack path will have larger crack deflection, higher fracture surface roughness and consequently slower crack growth rate ^{[29 , 30]}.