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j.matchar.2018.07.010

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Materials Characterization 144 (2018) 173–181
Contents lists available at ScienceDirect
Materials Characterization
journal homepage: www.elsevier.com/locate/matchar
Interface characteristics of Ti6Al4V-TiAl metal-intermetallic laminate (MIL)
composites prepared by a novel hot-pack rolling
T
⁎
Wei Suna,b,c, Fei Yangc, Fantao Konga,b, , Xiaopeng Wangb, Yuyong Chena,b
a
State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China
c
Waikato Centre for Advanced Materials, School of Engineering, The University of Waikato, Hamilton 3240, New Zealand
b
A R T I C LE I N FO
A B S T R A C T
Keywords:
Ti6Al4V/TiAl laminate composite
Hot-pack rolling
Interface microstructure
Phase constitution
Ti6Al4V (wt%)/Ti-44Al-8Nb-0.2W-0.2B-0.03Y (at.%) (high Nb-TiAl) metal–intermetallic laminate (MIL) composite with a defect-free interface was successfully prepared by hot-pack rolling and the interface microstructure
was characterized by XRD, SEM, EBSD, and TEM techniques. The results showed that the interface was about
260 μm thick and consisted of four different microstructure regions. The thickness of each region was about
60–80 μm and the microstructure of the four regions consisted of acicular O, α2, and β/B2 phases; acicular α2
phase; acicular α2 and β/B2 phases; and acicular γ, α2, and β/B2 phases respectively. The resulting microstructure of the interface region was attributed to the interdiffusion of Ti atoms from the starting Ti6Al4V alloy
layer to the high Nb-TiAl alloy layer and Al and Nb atoms from the high Nb-TiAl alloy layer to the Ti6Al4V alloy
layer. The formation of acicular α2 and γ phases was diffusion-controlled and the Nb segregation at the interface
caused the formation of O phase which had specific relationships with α2 and β/B2 phases. The detailed interface microstructural evolution of the Ti6Al4V/high Nb-TiAl MIL composite and the Vickers hardness of the
composite are also discussed in this paper.
1. Introduction
Ti-Al intermetallic alloys are attracting more and more attention
because of their low density, high specific modulus, and excellent oxidation resistance at high temperature [1–3]. However, the poor dislocation mobility and insufficient number of slip or twinning systems
lead to low toughness and ductility of the Ti-Al intermetallics, limiting
their extensive application.
Due to the high toughness, strength, and ductility of titanium alloy
[4,5], Ti/Ti-Al-based metal–intermetallic laminated (MIL) composites,
such as Ti/Al3Ti [6,7] and Ti/TiAl [8,9], are prepared to improve the
fracture toughness and ductility of the Ti-Al intermetallics at ambient
temperature. Various methods have been developed to prepare Ti/Ti-Al
MIL composites, such as diffusion bonding [10–12], welding [13,14],
and so on. However, micro-voids caused by interdiffusion or a reaction
between Ti and Al are formed at the interface between Ti-Al and Ti or in
the Ti-Al intermetallic layer of the Ti/Ti-Al MIL composites prepared by
the diffusion bonding approach and have the potential to decrease the
mechanical properties of the MIL composites. When using the welding
method, cracking is easily initiated in the interface region close to the
Ti-Al intermetallic layer, because high residual stress is likely to be
⁎
induced by the rapid cooling, and as a result, the mechanical properties
of the prepared MIL composites are promptly reduced.
Hot-pack rolling is an efficient way to manufacture metal–intermetallic MIL composites [15,16], in which the metal and intermetallic layers can be significantly deformed and then quickly stick
together, accelerating the atomic diffusion between them. Thus, the
formation of micro-voids at the interface between the metal and intermetallic layers can be avoided. The residual stress at the interface is
small due to the slow furnace cooling after the rolling, and as a result
there is less possibility of crack formation. Furthermore, a fine microstructure will be introduced in the MIL composites after hot-pack
rolling, which will not only improve the strength of the MIL composites
but also promote an interface reaction between the metal and the intermetallic. High Nb-TiAl alloys have much better high-temperature
oxide resistance than most TiAl alloys, even when the working temperature is over 800 °C [17,18]. Thus, Ti/high Nb-TiAl MIL composites
are very promising materials, with great potential to be used at higher
temperature than other Ti/TiAl MIL composites. The mechanical
properties of MIL composites are significantly dependent on the interface microstructure. However, few studies on the formation and constitution of the interface microstructure of Ti/TiAl MIL composites
Corresponding author at: State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China.
E-mail address: kft@hit.edu.cn (F. Kong).
https://doi.org/10.1016/j.matchar.2018.07.010
Received 27 January 2018; Received in revised form 27 April 2018; Accepted 6 July 2018
Available online 07 July 2018
1044-5803/ © 2018 Elsevier Inc. All rights reserved.
Materials Characterization 144 (2018) 173–181
W. Sun et al.
rolling speed was approximately 45 mm/s and the thickness reduction
was about 10–20% per pass. Between each rolling pass, the sample was
reheated up to 1250 °C and the temperature was held for 10 min. The
overall thickness reductions were about 50 and 60% of the original
thickness of the sandwich structure, respectively. After rolling at
1250 °C, the as-rolled sandwiches were transferred to the furnace for
annealing for 6 h at a preset temperature of 900 °C and then furnace
cooled down to room temperature. The MIL sheets were obtained by
machining off steel with dimensions of 170 mm × 70 mm × 3.2 mm for
the 50% reduction composite and 195 mm × 75 mm × 2.6 mm for the
60% reduction composite. The final one-layer thicknesses of the
Ti6Al4V and high Nb-TiAl alloys were 1.5 and 0.85 mm for the 50%
reduction composite and 1.3 and 0.65 mm for the 60% reduction
composite.
prepared by hot-pack rolling have been reported, with the exception of
research focused on the mechanical properties of hot-pack rolled Ti/
TiAl MIL composites [8,19].
In this paper, Ti6Al4V (wt%)/Ti-44Al-8Nb-0.2W-0.2B-0.03Y (at.%)
(high Nb-TiAl alloy) MIL composite was successfully produced by hotpack rolling at 1250 °C. The element diffusion behavior, phase composition and interfacial microstructure evolution was detailedly studied
in this paper, in an attempt to reveal the bonding mechanism of
Ti6Al4V and high Nb-TiAl alloy. To optimize the interfacial microstructure, the Vickers hardness and expected mechanical properties of
each interfacial region was also discussed.
2. Experimental Procedure
2.1. Preparation of Materials
2.3. Phase Analysis and Microstructure Characterization
The high Nb-TiAl alloy ingot was prepared from sponge Ti, highpurity Al, Al-Nb, Al-Y, and Al-W master alloys, and elemental boron
powders via vacuum arc remelting (VAR) and then the ingot was hot
isostatically pressed (HIPed) at 1300 °C for 4 h under a pressure of
130 MPa to reduce the shrinkage and small pores. Subsequently the
HIPed ingot was further homogenized at 950 °C for 36 h in air to reduce
and/or eliminate residual stress and composition segregation. After
that, the ingot was sealed in a stainless steel can 220 mm in diameter
and 500 mm in length and then forged at 1270 °C in one direction into a
pancake. The deformation strain rate was about 0.04 s−1 and the
overall deformation degree was about 70%. A schematic map of the
process described above is shown in Fig. 1. Four pieces of high Nb-TiAl
plates with dimensions of 100 mm × 65 mm × 2.2 mm were cut from
the as-forged high Nb-TiAl pancake and, together with two pieces of
Ti6Al4V plate with the same dimensions as the high Nb-TiAl plates,
were used to make two MIL composites through hot-pack rolling.
The phase constitutions and microstructure characteristics of the
produced Ti6Al4V/high Nb-TiAl MIL composite sheet were examined
by X-ray diffractometry (XRD), electron back scattered diffraction
(EBSD), scanning electron microscopy (SEM) in conjunction with energy dispersive X-ray spectrometry (EDS), and transmission electron
microscopy (TEM). The XRD was carried out using Cu Kα radiation
(λ = 0.154157 nm) with 2θ ranging from 20° to 100°. Specimens for
TEM observation were prepared by a standard procedure and an electrolytic jet polisher using electrolyte composed of 60% methanol, 30%
n-butyl alcohol, and 10% perchloric acid. The pore made by the electrolytic jet polisher covered the Ti alloy, interface, and TiAl alloy, and
TEM observation of the interface was carried out from the Ti alloy to
the TiAl alloy. The hardness measurements of the composite were made
with a standardized Vickers test device using a load of 500 g and a
dwell time of 15 s (at least five measurements).
3. Results
2.2. Hot-pack Rolling and Annealing
The cross-sectional microstructure (along the thickness direction) of
the 50% reduction Ti6Al4V/high Nb-TiAl MIL composite sheet produced by hot-pack rolling is shown in Fig. 3. It can be seen that strip
microstructures are formed after hot-pack rolling and these strips are in
different colors: grey, dark grey, and bright grey, corresponding to the
initial high Nb-TiAl, Ti6Al4V and the interface formed between the
high Nb-TiAl and Ti6Al4V plates, respectively. The interface has a
thickness of about 260 μm and no visible defects such as cracks or voids.
The Kirkendall pores reported in Van Loo [20] are successfully avoided
in the interface layer, as well. A high-magnification SEM image of the
interface, shown in Fig. 4a, reveals that it contains four microstructure
regions, each with a thickness of about 60–80 μm. Fig. 4b shows the
distribution of the chemical composition of Ti, Al, V, and Nb along the
line from position a of the Ti6Al4V alloy (region A) to position b of the
high Nb-TiAl alloy (region F) in Fig. 4a. The contents of Ti and V gradually decreased from position a to b, and the overall changes in the
contents of Al and Nb are opposite to those of Ti and V. This indicates
that atomic interdiffusion between the high Nb-TiAl and Ti6Al4V plates
took place during the processes of hot-pack rolling and annealing. A
sharp increase of Nb content at a position in region B is observed
(Fig. 4b), indicating that Nb-rich phase is probably formed in region B.
XRD analysis (Fig. 5) shows that the Ti6Al4V/high Nb-TiAl MIL composite is composed of α2-Ti3Al, α-Ti, β/B2, γ-TiAl, and O-Ti2AlNb
phases, of which O-Ti2AlNb phase is not supposed to be formed in either the high Nb-TiAl or the Ti6Al4V alloy. The possible formation
mechanism of O-Ti2AlNb phase in the MIL composite will be discussed
later.
EBSD analysis was used to identify the phase distribution across the
interface region, and the results are presented in Fig. 6. Four different
microstructure regions can be clearly observed in Fig. 6b, each composed of different phases: region B comprises O, α2, and β/B2 phases;
Alternating high Nb-TiAl and Ti6Al4V alloy plates were stacked to
form a ‘sandwich’ structure, with two high Nb-TiAl plates placing on
the two sides of the sandwich structure and a Ti6Al4V plate in the
middle. After that the ‘sandwiches’ were put in a stainless steel can
which was then vacuum pumped and sealed for hot-pack rolling; a
schematic map is shown in Fig. 2. The canned sandwiches were heated
to 1250 °C at a heating rate of 10 °C/min and then the temperature was
held for 1 h prior to hot rolling. The rolling was performed using a twohigh mill with rolls 350 mm in diameter and 300 mm in width. The
Fig. 1. The schematic diagram showing the process of producing high Nb-TiAl
ingot.
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W. Sun et al.
Fig. 2. Schematic map showing the MIL composite
sandwich stacking arrangement and the canning structure.
Fig. 5. XRD pattern of the 50% reduction Ti6Al4V/high Nb-TiAl MIL composite
sheet.
Fig. 3. SEM microstructure of the cross-sectional area of the as-rolled 50%
reduction Ti6Al4V/high Nb-TiAl MIL composite (through-thickness section).
15.41 at.%. This suggests that Ti2AlNb phase is formed in region B,
confirming the same result as was obtained from the EBSD and XRD
analyses. A needle-like structure is exhibited in region C (Fig. 7c), but it
is not discernable in the SEM image because of its fine microstructure.
More details on the needle-like microstructure will be presented in the
TEM image. Fig. 7d and f shows the microstructure of regions D and E; a
thin lamellar and needle-like structure can be clearly seen, and the
dimension of the lamellae in region E is finer than in region D. Equiaxed
precipitation and lamellar precipitation can be observed in Fig. 7f. The
EDS results show that the Nb concentration is 9.65 at.% in region D and
9.85 at.% in region E, which is slightly higher than the Nb concentration of 7.37 at.% found in region F. In addition, the Ti and Al contents
are 48.65 and 43.98 at.% in region F, respectively, indicating that the
chemical composition in area F is similar to that in the starting materials of high Nb-TiAl alloy.
Fig. 8 shows the TEM images of the interface area of 50% reduction
Ti6Al4V/high Nb-TiAl MIL composite which are taken from regions B
to E in Fig. 4, respectively. Acicular O and α2 phase can be clearly seen
in Fig. 8a. The dimensions of the α2 phase are a width of about 1 μm
region C comprises α2 phases; region D comprises α2 and β/B2 phases;
and region E comprises γ, α2 and β/B2 phases.
Fig. 7 shows high-magnification images of the Ti6Al4V alloy, high
Nb-TiAl, and the interface region in the Ti6Al4V/high Nb-TiAl MIL
composite, and the related EDS area analysis results of the different
areas shown in Fig. 4 are listed in Table 1.
In region A, the microstructure is mainly composed of thick lath α
phase and a small amount of β phase which is between the lath α phases
(Fig. 7a). EDS analysis shows that the Nb content is zero in this area,
and the concentrations of Al and V are identical to those of the starting
material of Ti6Al4V alloy. The mixed lamellae, with lengths of up to
50 μm, and equiaxed microstructure can be observed in Fig. 7b. Compared to region A, the Al concentration is significantly increased from
8.95 to 20.93 at.%; however, the content of Ti is decreased from 87.59
to 54.60 at.% and the V concentration is reduced from 3.46 to 0.77 at.
%. Nb also obviously appears in region B, with a concentration of
Fig. 4. (a) High magnification SEM image of the interface region in 50% reduction Ti6Al4V/high Nb-TiAl MIL composite, and (b) line scanning analysis of chemical
composition analysis of Ti, Al, V, and Nb along the line showed in (a).
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W. Sun et al.
Table 1
Chemical composition analysis for the different area in Fig. 4 and the starting
materials of high Nb-TiAl and Ti6Al4V alloy.
Position
Ti (at.%)
Al (at.%)
V (at.%)
Nb (at.%)
Possible phase
Region A
Region B
Region C
Region D
Region E
Region F
High Nb-TiAl
Ti6Al4V
87.59
61.60
68.16
54.78
50.34
48.65
48.42
87.71
8.95
22.22
27.38
35.57
39.81
43.98
43.94
8.73
3.46
0.77
0.96
0
0
0
0
3.56
0
15.41
3.50
9.65
9.85
7.37
7.64
0
α+β
α2 + β + O
α2
α2 + B2
α2 + γ + B2
α2 + γ + B2
α2 + γ + B2
α+β
TiAl MIL composite. However, due to the greater reduction in thickness,
the lath α phase translates into acicular α phase in the starting Ti6Al4V
alloy (Fig. 9b, region 1) and the γ and α2 phases in the high Nb-TiAl
alloy (Fig. 9c, region 6) are elongated.
As for the interfacial region, the lengths of the four regions at the
interface are all elongated, which means that the total thickness of the
interface of the 60% reduction composite (280 μm) is a little thicker
than that of the 50% reduction composite (260 μm). This may be because the 60% reduction composite is fabricated by seven rolling
passes, two more than were used for the 50% reduction composite.
Between each rolling pass, the composite is reheated up to 1250 °C and
held at that temperature for 10 min. The high temperature accelerates
the atomic diffusion and thus thickens the thickness of the interface.
Fig. 9d and e shows high-magnification SEM images of the interface of
the 60% reduction composite. They show that the phase composition of
each region at the interface of the 60% reduction composite is the same
as that of the 50% reduction composite. However, due to the greater
thickness reduction, the acicular α2 phase in regions 2, 3, and 4 (corresponding to regions C, D, and E in Fig. 4a) and the acicular γ phase in
region 4 are finer than in the 50% reduction composite. In conclusion,
the rolling process (including the rolling temperature, total thickness
reduction, rolling passes, and so on) may influence the microstructure
and the length of the interface but not determine the phase composition
of the composite.
Fig. 6. EBSD analysis result of the interfacial region of 50% reduction Ti6Al4V/
high Nb-TiAl MIL composite: (a) the front backscatter electron image; (b)
schematic of phase distribution.
and length of over 4 μm, and the acicular O is coarser than the acicular
α2 phase. In Fig. 8b, only α2 lamellar structure with a thickness of about
2 μm is exhibited. β/B2 and α2 phases are shown in Fig. 8c and the
thickness of the acicular α2 is up to 0.5 μm. Equiaxed β/B2, α2, and γ
phases and acicular γ phase are presented in Fig. 8d.
4. Discussion
4.1. Rolling Process
Fig. 9 shows the microstructure of the Ti6Al4V/high Nb-TiAl MIL
composite under seven rolling passes with a thickness reduction of
about 60%. Fig. 9a shows that after seven rolling passes, the composite
is also composed of six regions and the phase composition of each region is the same as in Fig. 4a, which indicates that the rolling process
will not determine the final phase composition of the Ti6Al4V/high Nb-
Fig. 7. High magnification SEM images of 50% reduction Ti6Al4V/high Nb-TiAl MIL composite at the different areas in Fig. 4: (a)–(f) represent areas A to F,
respectively, in Fig. 4.
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Materials Characterization 144 (2018) 173–181
W. Sun et al.
Fig. 8. Bright field TEM images of the microstructures of the interfacial region of 50% reduction Ti6Al4V/high Nb-TiAl MIL composite: the images of (a)–(d) are
taken from the region B to E in Fig. 4, respectively.
Fig. 9. The microstructure of the Ti6Al4V/high Nb-TiAl MIL composite with a thickness reduction about 60%: (a) the SEM image of the composite; (b)–(e) the high
magnification SEM images of each region corresponding to a.
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Materials Characterization 144 (2018) 173–181
W. Sun et al.
Fig. 10. Schematic diagram of microstructural evolution at the interface region between high Nb-TiAl and Ti6Al4V in the MIL composite: (a) atomic diffusion, (b)
Phase constitution just hot-pack rolling at 1250 °C, (c) the formation γ phase, (d) and (e) the formation of α2 phase and the migration of Nb atoms, (f) the segregation
of Nb atoms, and (g) the formation of O phase.
4.2.2. The Microstructural Evolution of Region E
The microstructural evolution at the interface between high Nb-TiAl
and Ti6Al4V in the MIL composite during the hot-pack rolling and
annealing is complicated. The schematic diagrams, as shown in Fig. 10,
illustrate how the four different microstructure regions are formed at
the interface between high Nb-TiAl and Ti6Al4V.
Overall, the atomic diffusion of Ti, Al, V, and Nb is activated at
1250 °C and there is a chemical concentration gradient between high
Nb-TiAl and Ti6Al4V (Fig. 4a). Thus, the Ti and V atoms diffuse from
Ti6Al4V alloy to high Nb-TiAl alloy, and the Al and Nb atoms diffuse in
the opposite direction to the Ti and V ones, that is, from high Nb-TiAl
alloy to Ti6Al4V alloy, as shown in Fig. 10a. Furthermore, the atomic
diffusion process is significantly accelerated by large plastic deformation induced by the hot-pack rolling because of the movement of dislocations and vacancies [24]. As a result, an interface layer is formed
between Ti6Al4V and high Nb-TiAl alloy after hot-pack rolling at
1250 °C. According to the Ti-Al-10Nb diagram (Fig. 11) [25] and chemical composition analysis of the different areas shown in Table 1,
when the temperature is up to 1250 °C, region E is in the α + β twophase region and is formed by β matrix and a little α phase.
After hot-packing rolling, the Ti6Al4V/high Nb-TiAl MIL composite
is transferred immediately to the furnace preset to 900 °C for annealing.
According to the Ti-Al-Nb ternary phase diagram (as shown in Fig. 12)
and the research by Koo et al. [26,27], region E, which contains 50 at.%
4.2. Annealing Process
4.2.1. The Starting High Nb-TiAl and Ti6Al4V Alloy
When the Ti6Al4V/high Nb-TiAl MIL composite is rolled at 1250 °C,
the Ti6Al4V layer is in the single β region [21] and the high Nb-TiAl
layer is in the α + γ region [22]. During rolling, the microstructures of
both Ti6Al4V and high Nb-TiAl are significantly broken due to the large
plastic deformation induced. When the MIL composite is annealed at
900 °C for 6 h immediately after hot rolling at 1250 °C, α phase is rapidly nucleated and grows in the Ti6Al4V alloy to form a lath structure.
No lamellar α/β lamellar structure is formed in the Ti6Al4V, because
the annealing temperature of 900 °C is very close to the α/β transfer
temperature of 995 °C [21] and holding the MIL composite at 900 °C for
6 h is sufficient to keep the transformation from β to α close to the nearequilibrium state. This leads to the formation of fine lath α and a very
small amount of β lamellae in the starting Ti6Al4V layer (region A in
Fig. 4). In the starting layer of high Nb-TiAl alloy, the α phase is quickly
ordered into α2, and β transfers into B2 phase during the transfer of the
as-rolled MIL composite to the furnace preheated to 900 °C and annealing. Then, γ phase precipitates from the α2 to form α2/γ lamellae
based on the well-developed theory [23]. This is the reason why the
microstructure of the high Nb-TiAl alloy in the as-rolled MIL composite
is composed of equiaxed γ, α2, and β/B2 phases and α2/γ lamellae.
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Materials Characterization 144 (2018) 173–181
W. Sun et al.
than region E, and therefore region D is in the β phase region during the
rolling process (Figs. 10b and 11). When the MIL composite sheet is
held at 900 °C for 6 h, α2 phase will precipitate from β phase. The phase
transformation in the interface region is near equilibrium, so β/α2
transformation is diffusion controlled rather than following the massive, bainite, or martensitic mechanism. As a result, the α2 phase nucleates and grows along the β {011} planes with the assistance of local
shear along the 〈110〉 directions and diffusion of atoms. The pole figure
in Fig. 13b shows that β phase and α2 phase have a fixed Burgers-type
orientation relationship: {110}β // {0001}α2 and {111}β // {112 0}α2,
confirming that α2 precipitates from the β phase and the {110} planes
of β phase are the habit planes of α2 phase. Furthermore, the α2 phase
grows stably along the β {110} plane, because the habit planes always
have low-energy boundaries. As a result, α2 phase grows preferentially
in the length direction rather than in the thickness direction, leading to
the formation of acicular α2 phase by the lateral movement of small
ledges or steps along the coherent plates [30]. Meanwhile, the Nb atoms
are expelled from the α2 phase in region D and diffuse into the adjacent
β-rich area, region C, as illustrated in Fig. 10d. This is because Nb is a
strong β phase stabilizer and has lower solubility in α2 phase than in β
phase, and the growth of α2 phase requires the consumption of Al atoms
during the β/α2 transformation and promotes diffusion of the solute Al
towards α2 phase and diffusion of Nb in the opposite direction to the
retained β phase [31].
Fig. 11. Quasi-phase diagrams showing effect of 10% Nb additions on phase
relationship of near γ-TiAl alloys.
Ti, 40 at.% Al, and 10 at.% Nb, will reach the phase region of α + γ + β
when the temperature is 1250–1200 °C (α phase will also order into α2
phase at 1200 °C [28]). Thus, γ phase may precipitate from β phase in
the interface region E, which is close to the starting high Nb-TiAl layer,
during the process of transferring the as-rolled MIL composite to the
furnace preset to 900 °C and annealing at 900 °C. However, no α2/γ
lamellae are observed in the interface region E (as shown in Figs. 6b
and 8d), suggesting that most of the γ phase is precipitated in β matrix
to form acicular γ and equiaxial γ, a little γ phase is precipitated from
α2 phase to form acicular γ and equiaxial γ, and no γ phase is precipitated from α2 phase to form α/γ or α2/γ lamellae. This is attributed
to two factors [26]: (1) the vacancy volume diffusion rate in β phase is
faster than that in α2 phase, and consequently atomic consumption for
the nucleation and growth of γ phase is easier in the β phase; (2) the
characteristics of the coherent or semi-coherent α/γ or α2/γ interface
make it difficult for γ to grow in the α phase. Fig. 13a suggests that γ
phase and α2 phase exist a fixed Blackburn-type orientation relationship: namely {111}γ // {0001}α2 and {110}γ // {112 0}α2, and the γ
phase and β phase have the orientation relationship {111}γ // {110}β
and {011}γ // {111}β. This partially confirms that γ phase nucleates
from the β phase, and similar results have also been reported by other
researchers [29–32].
4.2.4. The Microstructure Evolution of Region C
Section 4.2.3 mentions that the α2/β transformation temperature
will be significantly increased with increases of the Al concentration;
thus region C is also in the β phase region during the rolling process as
the Al concentration in region C is lower than that in region D. When it
comes to the annealing process, α2 phase grows preferentially in the
region with lower Al content and its volume fraction is increased with
the decrease of Al content [32]; thus the amount and size of α2 phase
will be increased compared with region D below 1150 °C [25], eventually leading to a large amount of α2 phase with relatively large dimensions, and a very small amount of equiaxed β/B2 phase (the β
phase will order into B2 phase when the temperature is below 1100 °C
[27]) is observed in region C, as illustrated in Fig. 10e, f, and g. With
prolongation of the annealing time at 900 °C, the α2 phase continues to
consume Al solute and grow, while the Nb solute continues to be expelled towards region B; as a result, no obvious β/B2 phase is observed,
while a large amount of α2 phase exists in region C, as shown in Figs. 6
and 7.
4.2.3. The Microstructure Evolution of Region D
Al is a strong α phase stabilizer, and the α2/β transformation temperature will be significantly increased with increases of the Al concentration [29]. Table 1 shows that region D has a lower content of Al
4.2.5. The Microstructure Evolution of Region B
Like regions C and D, region B is also in the β phase region during
the rolling process and the α2 phase will precipitate from the β phase
during the annealing process. However, as discussed in Section 4.2.1, a
Fig. 12. Isothermal phase diagrams of the Ti-Al-Nb ternary system at 1200 °C and 1300 °C.
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Materials Characterization 144 (2018) 173–181
W. Sun et al.
Fig. 13. The EBSD analysis of the interfacial region: (a) region E and the pole figure of β/B2 and γ phase; (b) region D and the pole figure of β/B2 and α2 phase; (c)
region B and the pole figure of β/B2, α2 and O phase (α2 in red color, β/B2 in yellow color, γ in blue color and O in pink color). (For interpretation of the references to
color in this figure legend, the reader is referred to the web version of this article.)
large amount of α phase is formed in the starting Ti6Al4V layer (region
A) during the process of annealing at 900 °C, so it is difficult for the Nb
atoms in region B to further dissolve into the region A because of the
limited solubility of Nb in the α phase [31], and this is confirmed by the
EDS analysis (Table 1) showing the zero content of Nb in region A. As a
result, a large amount of Nb is accumulated in region B, with a relatively high value of 15.41 at.%.
According to the ternary phase diagram of Ti-22Al-xNb (Fig. 14)
[33], the O phase starts to precipitate from β/B2 phase when the
temperature is below 950 °C. In region B, the phase transformation will
take place along the red line indicated in Fig. 14, and the final phase
constitution is composed of α2, β/B2, and O, having a lamellar structure. The pole figure (Fig. 13c) shows that O phase and α2 phase have a
fixed orientation relationship: {001}O // {0001}α2 and {010}O //
{1100}α2, and O phase and β/B2 phase have an orientation
relationship: {001}O // {110}β/B2 and {010}O // {110}β/B2. This
further confirms the precipitation of O phase from β/B2 and the existence of three phases in region B in the final microstructure of the
interface region of the MIL composite.
4.3. The Vickers Hardness of the Composite
Fig. 15 shows the Vickers hardness of the Ti6Al4V/high Nb-TiAl
MIL composite in different regions. From Ti6Al4V to high Nb-TiAl
alloy, the Vickers hardness shows a trend of an initial decrease followed
by an increase and finally a decrease. Region E has the highest hardness
of about 456.56 Hv and region C has the lowest hardness of about
282.14 Hv. It is well known that the hardness and strength are always
positively correlated and the hardness and toughness are always negatively correlated. So, regions B and C may have better toughness at
room temperature and regions D and E may have higher strength at
room temperature.
Fig. 15. Vickers hardness of the 50% reduction Ti6Al4V/high Nb-TiAl MIL
composite. The value in horizontal ordinate corresponds to the six regions in
Fig. 4. Error bars represent the standard deviation during measurements.
Fig. 14. Quasi-binary sections of the equilibrium phase diagram of
Ti–22Al–xNb (at.%) phase diagram.
180
Materials Characterization 144 (2018) 173–181
W. Sun et al.
Previous study [34] shows that the ductility and fracture toughness
of Ti2AlNb alloy at room temperature are much higher than those of
TiAl alloy and the ultimate tensile strength of Ti2AlNb alloy can also
exceed that of the TiAl alloy. Thus region B is beneficial for the mechanical properties of Ti/TiAl MIL composite as it may not only provide
higher ductility and toughness but also maintain the high strength at
room temperature. Compared with region B, region C has lower hardness, which indicates that region C may have better toughness at room
temperature.
As for regions D and E, the higher content of B2 phase increases the
hardness of these two regions, and may thus increase the strength of
these two regions and the bonding strength between Ti6Al4V and high
Nb-TiAl alloy. As we know, for MIL composite, the interface plays an
important role in improving the toughness and plasticity. Weak interfacial bonding will lead to preferential crack initiation at interfaces,
which may harm the mechanical properties of the MIL composites. The
higher strength of regions D and E may ensure a high level of load
transfer and thus increase the mechanical properties of the Ti6Al4V/
high Nb-TiAl MIL composites.
[7]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
[15]
5. Conclusions
[16]
(1) Ti6Al4V (wt%)/Ti-44Al-8Nb-0.2W-0.2B-0.03Y (at.%) (high NbTiAl) MIL composite was successfully produced by hot-pack rolling
at 1250 °C, and there were no pores or cracks in the interface between high Nb-TiAl and Ti6Al4V.
(2) Four microstructure regions, with an overall thickness of about
260 μm, were formed in the interface between high Nb-TiAl and
Ti6Al4V. They were composed of acicular α2, O, and β/B2 phases;
lamellar α2 phase; acicular α2 and β/B2 phases; and acicular γ,
equiaxed α2, and β/B2 phases, respectively, from the Ti6Al4V to
the high Nb-TiAl side.
(3) The formation of acicular α2 and γ phases in the interface region is
primarily controlled by diffusion, and the formation of O phase in
the region close to the Ti6Al4V starting layer is caused by the
segregation of Nb content.
(4) The Vickers hardness of the composite shows a trend of an initial
decrease followed by an increase and finally a decrease. Region E
has the highest hardness of about 456.56 Hv and region C has the
lowest hardness of about 282.14 Hv.
[17]
[18]
[19]
[20]
[21]
[22]
[23]
[24]
[25]
[26]
Acknowledgement
[27]
This work was financially supported by the National Natural Science
Foundation of China (Project NO. 51471056).
[28]
[29]
References
[30]
[1] Y.W. Kim, Ordered intermetallic alloys, part III: gamma titanium aluminides, JOM
46 (1994) 30–39.
[2] M. Yamaguchi, Y. Umakoshi, The deformation-behavior of intermetallic superlattice compounds, Prog. Mater. Sci. 34 (1990) 1–148.
[3] F. Appel, R. Wagner, Microstructure and deformation of two-phase gamma-titanium
aluminides, Mater. Sci. Eng. R 22 (1998) 187–268.
[4] H.L. Yu, M. Yan, J.T. Li, A. Godbole, C. Lu, K. Tieu, H.J. Li, C. Kong, Mechanical
properties and microstructure of a Ti-6Al-4V alloy subjected to cold rolling,
asymmetric rolling and asymmetric cryorolling, Mater. Sci. Eng. A 710 (2018)
10–16.
[5] F. Yang, B. Gabbitas, Feasibility of producing Ti-6Al-4V alloy for engineering application by powder compact extrusion of blended elemental powder mixtures, J.
Alloys Compd. 695 (2017) 1455–1461.
[6] H.L. Yu, C. Lu, A.K. Tieu, H.J. Li, A. Godbole, C. Kong, Annealing effect on
[31]
[32]
[33]
[34]
181
microstructure and mechanical properties of Al/Ti/Al laminate sheets, Mater. Sci.
Eng. A 660 (2016) 195–204.
P.J. Zhou, C.H. Guo, E.H. Wang, Z.M. Wang, Y. Chen, F.C. Jiang, Interface tensile
and fracture behavior of the Ti/Al3Ti metal-intermetallic laminate (MIL) composite
under quasi-static and high strain rates, Mater. Sci. Eng. A 665 (2016) 66–75.
F.T. Kong, Y.Y. Chen, Preparation of γ-TiAl/TC4 composite sheet and its microstructure and properties, Rare Metal Mater. Eng. 38 (2009) 1484–1486.
H. Li, C. Yang, L.X. Sun, M.Q. Li, Hot press bonding of gamma-TiAl and TC17 at a
low bonding temperature by imposing plastic deformation and post heating, Mater.
Lett. 187 (2017) 4–6.
X.R. Wang, Y.Q. Yang, X. Luo, W. Zhang, G.M. Zhao, B. Huang, An investigation of
Ti-43Al-9V/Ti-6Al-4V interface by diffusion bonding, Intermetallics 36 (2013)
127–132.
Y.Q. Han, C.F. Lin, X.X. Han, Y.P. Chang, C.H. Guo, F.C. Jiang, Fabrication, interfacial characterization and mechanical properties of continuous Al2O3 ceramic fiber
reinforced Ti/Al3Ti metal-intermetallic laminated (CCFR-MIL) composite, Mater.
Sci. Eng. A 688 (2017) 338–345.
E.H. Wang, Y. Tian, Z.Q. Wang, F.F. Jiao, C.H. Guo, F.C. Jiang, A study of shape
memory alloy NiTi fiber/plate reinforced (SMAFR/SMAPR) Ti-Al laminated composites, J. Alloys Compd. 696 (2017) 1059–1066.
K. Han, H.Q. Wang, B.G. Zhang, Y.X. Li, T. Wang, Effect of thermal compensation on
microstructure and mechanical properties of electron-beam welded joint for highNb containing TiAl/Ti600 alloys, Mater. Des. 131 (2017) 273–285.
G.Q. Chen, B.G. Zhang, W. Liu, J.C. Feng, Influence of aluminum content on the
microstructure and properties of electron beam welded joints of TiAl/TC4 alloy,
Rare Metal Mater. Eng. 42 (2013) 452–456.
J.G. Luo, V.L. Acoff, Using cold roll bonding and annealing to process Ti/Al multilayered composites from elemental foils, Mater. Sci. Eng. A 379 (2004) 164–172.
Y.M. Hwang, H.H. Hsu, Y.L. Hwang, Analytical and experimental study on bonding
behavior at the roll gap during complex rolling of sandwich sheets, Int. J. Mech. Sci.
42 (2000) 2417–2437.
Y.J. Su, F.T. Kong, Y.Y. Chen, N. Gao, D.L. Zhang, Microstructure and mechanical
properties of large size Ti-43Al-9V-0.2Y alloy pancake produced by pack-forging,
Intermetallics 34 (2013) 29–34.
Y.B. Zhao, S.Z. Zhang, C.J. Zhang, P. Lin, Z.P. Hou, Y.Y. Chen, Microstructural
evolution of hot-forged high Nb containing TiAl alloy during high temperature
tension, Mater. Sci. Eng. A 678 (2016) 116–121.
F.T. Kong, Y.Y. Chen, D.L. Zhang, Interfacial microstructure and shear strength of
Ti–6Al–4V/TiAl laminate composite sheet fabricated by hot packed rolling, Mater.
Des. 32 (2011) 3167–3172.
F.J.J. van Loo, G.D. Rieck, Diffusion in the titanium-aluminium system—II.
Interdiffusion in the composition range between 25 and 100 at.% Ti, Acta Mater. 21
(1973) 73–84.
C.F. Yolton, F.H. Froes, R.F. Malone, Alloying element effects in metastable betatitanium alloys, Metall. Trans. A. 10 (1979) 132–134.
V.T. Witusiewicz, A.A. Bondar, U. Hecht, T.Y. Velikanova, The Al–B–Nb–Ti system,
J. Alloys Compd. 472 (2009) 133–161.
Y.W. Kim, D.M. Dimiduk, Progress in the understanding of gamma titanium aluminides, JOM 43 (1991) 40–47.
N. Orhan, M. Aksoy, M. Eroglu, A new model for diffusion bonding and its application to duplex alloys, Mater. Sci. Eng. A 271 (1999) 458–468.
Y.F. Liang, X.J. Xu, J.P. Lin, Advances in phase relationship for high Nb-containing
TiAl alloys, Rare Metals 35 (2016) 15–25.
T.H. Yu, C.H. Koo, Microstructural evolution of a hot-rolled Ti–40Al–10Nb alloy,
Mater. Sci. Eng. A 239-240 (1997) 694–701.
R. Kainuma, Y. Fujita, H. Mitsui, I. Ohnuma, K. Ishida, Phase equilibria among α
(hcp), β (bcc) and γ (L10) phases in Ti–Al base ternary alloys, Intermetallics 8
(2000) 855–867.
X.H. Wu, D. Hu, M.H. Loretto, Alloy and process development of TiAl, J. Mater. Sci.
39 (2004) 3935–3940.
D. Banerjee, J.C. Williams, Perspectives on titanium science and technology, Acta
Mater. 61 (2013) 844–879.
K.L. Yang, J.C. Huang, Y.N. Wang, Phase transformation in the β phase of super α2
Ti3Al base alloys during static annealing and superplastic deformation at
700–1000°C, Acta Mater. 51 (2003) 2577–2594.
G.H. Liu, Z.D. Wang, T.L. Fu, Y. Li, H.T. Liu, T.R. Li, M.N. Gong, G.D. Wang, Study
on the microstructure, phase transition and hardness for the TiAl–Nb alloy design
during directional solidification, J. Alloys Compd. 650 (2015) 45–52.
Z.C. Liu, J.P. Lin, S.J. Li, G.L. Chen, Effects of Nb and Al on the microstructures and
mechanical properties of high Nb containing TiAl base alloys, Intermetallics 10
(2002) 653–659.
K. Muraleedharan, D. Banerjee, S. Banerjee, S. Lele, The α2-to-O transformation in
Ti-Al-Nb alloys, Philos. Mag. 71 (1995) 1011–1036.
J. Kumpfert, Intermetallic alloys based on orthorhombic titanium aluminide, Adv.
Eng. Mater. 3 (2001) 851–864.
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