Microstructural Characterization and Phase Development at the
Interface Between Aluminum Nitride and Titanium After Annealing at
13001–15001C
Chia-Hsiang Chiu and Chien-Cheng Lin*
,wDepartment of Materials Science and Engineering, National Chiao Tung University, Hsinchu 30050, Taiwan Diffusion couples of aluminum nitride (AlN) and Ti were
annealed under an argon atmosphere at temperatures ranging from 13001 to 15001C for 0.5–36 h. The morphologies, crystal structures, and chemical compositions of the reaction zones at AlN/Ti interfaces were characterized using analytical scanning electron microscopy and analytical transmission electron microscopy. An interfacial reaction zone, consisting of TiN, s2-Ti2AlN,s1-Ti3AlN,a2-Ti3Al, and a two-phase (a2-Ti3Al1a-Ti) region in sequence, was observed in between AlN and Ti after annealing at 13001C. Thea2-Ti3Al region revealed equiaxed and elongated morphologies with ½0001equiaxed j j ½1100elongated and ð1010Þequiaxed j jð1122Þelongated. In the two-phase (a2-Ti3Al1a-Ti) region, a2-Ti3Al and a-Ti were found to satisfy the follow-ing orientation relationship: ½0001a-Ti
j j½0001
Ti3Al and
ð1100Þa-Ti j j ð1100ÞTi3Al. The c-TiAl and a lamellar two-phase
(c-TiAl1a2-Ti3Al) structure, instead of s1-Ti3AlN, were found in between s2-Ti2AlN and a2-Ti3Al after annealing at 14001C. The orientation relationship of c-TiAl and a2-Ti3Al in the lamellar structure was identified to be as follows: ½011TiAl j j½2110Ti3Al and ð111ÞTiAl
j jð010Þ
Ti3Al. Compared
with the reaction zone after annealing at 14001C, the c-TiAl was not found at the interface after annealing at 15001C. The microstructural development resulting from isothermal diffusion at 13001C and subsequent cooling at the interface are explained with the aid of the Ti–Al–N ternary phase diagram and a mod-ified Ti–Al binary phase diagram.
I. Introduction
A
LUMINUM NITRIDE has been considered one of the most promising substrate materials for use in semiconductor, microwave, optic, electronic, and other high-performance appli-cations, because of its high thermal conductivity ( 320 W (m L)1), low dielectric constant ( 8.8), a high electric re-sistivity ( 1012–1014O cm), and coefficient of thermal expan-sion similar to that of silicon.1,2Meanwhile, Ti is a highly active metal and easily reacts with almost all ceramics,3–9and a high adhesion or bond strength can be achieved by the interface re-actions between titanium and the substrate materials. In some applications for microelectronics such as ceramic packaging and metallization,10–14 where aluminum nitride is placed in direct contact with titanium, the interfacial properties of AlN and Ti are crucial in determining the quality and reliability of the elec-tric package and its high-temperature applications.Previous investigations on the interfacial reaction of AlN and Ti have been very limited, most of them concerning the reaction
of AlN with Ti thin films15–19or Ti-containing brazing foils.20–22
While available data on the microstructure of the AlN–Ti in-terfacial reaction zone are often controversial, it is imperative to explore the interdiffusion reactions and mechanisms in the AlN/ Ti bonding process in order to make more effective use of AlN and Ti as well.
Some previous studies regarding the thin-film metallization of AlN have been published.15–19As far as the interface reaction between the Ti thin film and the AlN substrate was concerned, Westwood and Notis15,16found the formation of TiN and TiAl3
at the interface after annealing at 6001C for 30 min in an oxygen-free sample. He et al.23investigated the interface reac-tions of Ti thin films with AlN substrate in the temperature range of 6001–8001C using X-ray diffractometry (XRD) and Ruthford backscattering spectroscopy (RBS). They indicated that the TiAl3phase was formed at the interface adjacent to the
AlN substrate, while TiN, Ti4N3x, and Ti2N were formed
above the TiAl3layer. Yasurnoto et al.18 deposited a Ti thin
film on AlN with radio frequency (rf) sputtering, revealing that under an argon atmosphere, TiAl3 was formed at 7001C for
60 min, and TiAl3, Ti2N, and TiN were detected after annealing
at 8301C for 60 min by using XRD. In the study of the inter-diffusion and reaction of Ti (thin film) and AlN (substrate) using RBS and transmission electron microscopy (TEM), Iman-aka and Notis19found Ti2AlN at the interface after annealing
at 8001–9501C. Recently, Pinkas et al.24 worked on the early stages of interface reactions between Al and Ti thin films after annealing at 6001C for 1–10 h. They claimed that the AlN decomposed at the AlN/Ti interface and its products, Al and N, reacted with Ti to produce an AlN/Al3Ti/Ti2N/Ti3Al/a-(Ti,Al)ss
phase sequence.
As for the brazing of AlN,20–22 Carim and Loehman20 re-ported that continuous TiN and (Ti, Cu, Al)6N at the interface
of AlN and Ag–Cu–Ti foil were formed after annealing at 9001C for 5–30 min. By using TEM and electron probe microanalysis, Loehman21and Loehman and Tomsia22indicated that TiN
0.7
was detected at the interface of AlN and Ag–26.7Cu–4.5Ti after reaction at 9001C for 30 min in an argon atmosphere.
Among other previous studies on the interfacial reaction of AlN and Ti, El-Sayed et al.25 characterized the reaction zone microstructure of AlN/Ti (20 or 50 mm thick)/AlN joints after annealing at 10501–12001C for 2–20 h in vacuum. The reflection peaks of TiN, Ti2AlN, Ti3AlN, and Ti3Al were observed in the
X-ray diffraction spectra taken from the fracture surfaces of annealed joints. Up to 12001C for 20 h, both TiN1x and
Ti3AlN did not grow significantly, but the growth kinetics of
Ti3Al followed the parabolic law. Paransky et al.26–28
investi-gated the interfacial reactions between AlN particles and the Ti matrix, as well as AlN–Ti diffusion couples, after annealing in the temperature range from 9001 to 11001C using energy-dis-persive spectroscopy (EDS) and electron back-scattered diffrac-tion attached to a scanning electron microscope (SEM). A phase sequence of TiN, Ti3Al0.8N0.8, and Ti3Al was observed
at the interface of Ti and AlN. While the binary nitride TiN and the ternary nitride Ti3AlN exhibited a complex
inter-penetrating morphology, a lamellar two-phase region was also 1409
Journal
r2005 The American Ceramic SocietyR. Cutler—contributing editor
This work was supported by the National Science Council under the contract No. NSC 93-2216-E-009-017.
*Member, American Ceramic Society.
w
Author to whom correspondence should be addressed. e-mail: chienlin@faculty. nctu.edu.tw
observed between Ti3AlN and Ti3Al layers after annealing at
10001 and 11001C.
Many applications of the industrial AlN/Ti joints, such as metallization, brazing, and the composites mentioned above, are determined by the characteristics of the interface between AlN and Ti. In the last few decades, extensive studies have been car-ried out on the interface reaction between aluminum nitride and titanium. However, the microstructure evolution at the interface has not been elucidated to date, even though a fundamental understanding of reaction and diffusion mechanisms is of great importance for industrial applications and scientific meaning.
The present study is devoted to the microstructural charac-terization of the interfacial reaction zone in AlN and Ti diffusion couples after annealing at temperatures ranging from 13001 to 15001C by SEM/EDS and TEM/EDS. We will try to explain the microstructural development at the Ti/AlN inter-face on the basis of the ternary Al–Ti–N phase diagram and the diffusion paths that connect the phases formed by the reac-tion between AlN and Ti. The present study is expected to contribute to the understanding of the ternary Ti–Al–N system at high temperatures and to aid in the processing of ceramic– metal joints.
Table I. New Phases Formed in the Interfacial Reaction Zone of AlN/Ti Diffusion Couples
Annealing conditions
Phases
TiN Ti2AlN Ti3AlN TiAl
Lamellar structure
(TiAl1Ti3Al) Ti3Al
Two phase region (Ti3Al1Ti) 13001C/0.5 h
x x x 13001C/3 h x x 13001C/10 h x x 13001C/36 h x x 14001C/0.5 h x x x 14001C/3 h x 14001C/10 h x 14001C/36 h x 15001C/0.5 h x 15001C/3 h x x 15001C/10 h x x 15001C/36 h x x , observed; x, none. TiAl Ti2AlN TiN Ti Ti3Al AlN Ti3AlN Ti3Al+Ti (a)A
B
AlN Ti2AlN Ti3Al Ti3Al+Ti Ti TiN TiAl+Ti3Al (b)A
B
TiN Ti Ti3Al AlN Ti3Al+Ti Ti2AlN TiAl+Ti3Al (c)A
B
Fig. 1. Scanning electron micrographs of the interface between AlN and Ti after annealing at: (a) 13001C/36 h; (b) 14001C/36 h; and (c) 15001C/36 h.
II. Experimental Procedure
Highly pure AlN plates (SH-15, Tokuyama Soda Corp., Tokyo, Japan) and Cp–Ti plates (99.7% purity, Alfa Aesar, Ward Hill, MA) were used in this study. All the plates (about 15 mm 10 mm 4 mm in dimension) were ground with a diamond (15 mm) matted disk and then polished with a diamond paste (3 mm) and an alumina suspension (1 mm) using a precision polishing ma-chine (Model Minimet 1000, Buehler Ltd., Lake Bluff, IL). The specimens were then rinsed in acetone (ultrasonic bath) and dis-tilled water, and then air-dried.
In order to characterize the microstructure of the Ti/AlN in-terface, samples were prepared as a sandwich mode with the Ti metal placed in between two pieces of AlN. Then, the samples were annealed in an argon atmosphere (with O2o1 ppm, H2Oo
76 ppm, THCo0.5 ppm, and N2o3 ppm) under a pressure of 2
MPa at temperatures ranging from 13001 to 15001C, with a hold-ing time from 0.5 to 36 h, and then the specimen was continuously cooled down to room temperature at a rate of 101C/min.
The cross-sectional TEM and SEM samples were prepared as follows: each annealed sample was cut into two halves in a di-rection perpendicular to the interface of AlN/Ti, and then
ground and polished by the standard procedures as mentioned above. The samples were etched with the Kroll reagent (10 mL HF130 mL HNO3160 mL H2O) in order to emphasize the
features of different phases and to remove the deformed surface layer. Thereafter, the samples were rinsed in acetone (ultrasonic bath) and distilled water, and then air dried. To avoid charging, all the SEM samples were coated with a thin layer of platinum (finished preparation of SEM samples). The cross-sectional slab was ground down to an 80–100 mm thickness using a precision polishing machine; then, the sample was thinned to 20–30 mm by dimpling, and finally argon ion milled at 5 kV and 20 mA (fin-ished preparation of TEM samples).
Microstructural characterization of the cross-sections of AlN/ Ti was carried out using a high-resolution scanning electron microscope (Model JSM-6500, JEOL, Tokyo, Japan) and an analytical TEM (Model 2000Fx, JEOL). The Cliff–Lorimer standardless technique was used to analyze the compositions of the various phases. The technique was performed on the TEM, equipped with an ultra-thin window EDS detector (Mod-el 9900, EDAX International, Prairie View, IL). A conventional ZAF correction procedure included in the LINK ISIS software was used for the quantitative analyses.
Fig. 2. (a) Bright-field image of the AlN/Ti interface after annealing at 13001C for 3 h; (b) selected area diffraction pattern (SADP) of TiN, Z 5 [001]; (c) SADP of t2-Ti2AlN, Z¼ ½1120; and (d) SADP of t1-Ti3AlN, Z 5 [001].
III. Results and Discussion
The products formed in the AlN/Ti interfacial reaction zone are listed in Table I, after annealing at 13001–15001C for various periods. The reaction zone consisted of TiN, t2-Ti2AlN, t1
-Ti3AlN, a2-Ti3Al, and a two-phase (a2-Ti3Al1a-Ti) region in
sequence after annealing at 13001C. The g-TiAl and a lamellar two-phase (g-TiAl1a2-Ti3Al) structure were found instead of t1
-Ti3AlN in between t2-Ti2AlN and a2-Ti3Al after annealing at
14001C. In comparison with the results after annealing at 14001C, g-TiAl was not formed at the interface after annealing at 15001C. It was noted that there were some exceptions for the initial transition stage. For instance, no t2-Ti2AlN had been
found after annealing at 13501C for 0.5 h, while g-TiAl and the lamellar two-phase (g-TiAl1a2-Ti3Al) layer did not exist after
annealing at 14001C for 0.5 h, whereas a layer of g-TiAl was found after annealing at 15001C for 0.5 h. In contrast to the re-action of AlN and a-Ti thin films annealed at lower temperatures (e.g., 6001–8001C),23,24no TiAl
3was found in the present study.
(1) SEM/EDS Analyses
Figure 1 shows the secondary electron images of the cross-sectional microstructures of the AlN/Ti diffusion couples after
annealing at various temperatures for 36 h. Figure 1(a) displays that the reaction zone between AlN and Ti consists of TiN, t2
-Ti2AlN, t1-Ti3AlN, a2-Ti3Al, and a two-phase (a2-Ti3Al1a-Ti)
region in sequence after annealing at 13001C for 36 h. The microstructure of AlN/Ti after annealing at 14001C for 36 h is demonstrated in Fig. 1(b), which displays the reaction phase se-quence of TiN/t2–Ti2AlN/g-TiAl/(g-TiAl1a2-Ti3Al)/a2-Ti3Al/
(a2-Ti3Al1a-Ti). The g-TiAl layer and lamellar structure
(g-TiAl1a2-Ti3Al), instead of Ti3AlN, were formed between
the Ti2AlN and Ti3Al. Figure 1(c) displays the reaction phase
sequence of TiN/t2–Ti2AlN/(g-TiAl1a2-Ti3Al)/a2-Ti3Al/
(a2-Ti3Al1a-Ti). The TiAl layer was not found at the interface
after annealing at 15001C for 36 h as mentioned above. El-Sayed et al.25found that the sequence of the reaction layer was AlN/TiN/t1-Ti3AlN/a2-Ti3Al/a-Ti. However, they did not
observe the lamellar two-phase (g-TiAl1a2-Ti3Al) and (a2
-Ti3Al1a-Ti) layers. Paransky et al.26–28 showed that the
se-quence of TiN, t1-Ti3AlN, and a2-Ti3Al was formed at the
in-terface of Ti (matrix) and AlN (particle) and that the TiN and t1-Ti3AlN exhibited a complex interpenetrating morphology.
Furthermore, a lamellar two-phase region was also observed between t1-Ti3AlN and a2-Ti3Al layers after annealing at 10001
and 11001C. However, in the present study, t2-Ti2AlN and a
Fig. 3. After annealing at 13001C for 3 h: (a) bright-field (BF) image of the equiaxed a2-Ti3Al (abutting the t1-Ti3AlN layer in Fig. 1(a)); (b) BF image
of the elongated a2-Ti3Al (abutting the two-phase (a2-Ti3Al1a-Ti) layer in Fig. 1(a)); (c) variation in the grain size of the equiaxed a2-Ti3Al along
the direction from the AlN side (bottom) to the Ti side (top); (d) selected area diffraction pattern of both elongated and equiaxed a2-Ti3Al, showing the
orientation relationship of½0001equiaxed j j½1100elongatedandð1010Þequiaxed j jð1122Þelongated.
two-phase (a2-Ti3Al1a-Ti) region existed at the interface after
annealing at 13001C, while g-TiAl and a lamellar structure (g-TiAl1a2-Ti3Al) were found in the interface at 14001C. In the
temperature range used in this study, two lamellar layers were found: one layer (a-Ti1a2-Ti3Al) was developed because of the
precipitation of a2-Ti3Al from a-Ti; the other layer (a2-Ti3
Al1g-TiAl) was formed because of the eutectoid reaction ða-Ti ! a2-Ti3Alþ g-TiAlÞ during cooling.
In Fig. 1, the nitride layers, e.g., TiN, t2-Ti2AlN, and t1
-Ti3AlN, were so brittle that there were cracks in these layers.
The crack formation was attributed to the mismatch of thermal expansion coefficients between aluminum nitride and titanium, resulting in a significant residual thermal stress.
It was also noted that there existed some pores in aluminide layers (i.e., g-TiAl, a2-Ti3Al, etc.) of the reaction zone. These
pores were caused because of the formation of nitrogen bubbles during cooling. A significant amount of nitrogen was dissolved in aluminide layers on heating. However, the solubility of ni-trogen in aluminide layers sharply decreased with decreasing temperature, so that nitrogen was supersaturated. The nitrogen bubbles were thus precipitated through a nucleation and growth mechanism like the gas bubble formation commonly encoun-tered in the casting of alloys. It is worth noting that the decrease in the solubility of nitrogen, not the solubility itself, gave rise to
the precipitation of nitrogen bubbles or pores during cooling. This is the reason why nitrogen bubbles primarily existed in a2
-Ti3Al, although the solubility of nitrogen in a2-Ti3Al is much
larger than that in g-TiAl. Like the existence of spherical oxygen bubbles in Ti because of the Ti/ZrO2interfacial reaction,29
ni-trogen bubbles were formed in a-Ti even though the a-Ti is ca-pable of dissolving a large amount of oxygen. On comparing Figs. 1(a)–(c), it was concluded that the amount of bubbles in-creased with annealing temperature.
(2) TEM/EDS Analyses
Figure 2(a) shows the bright-field (BF) image of the AlN/Ti in-terface after annealing at 13001C for 3 h, showing the phases of TiN, t2-Ti2AlN, t1-Ti3AlN, and a2-Ti3Al. Note that the
two-phase (a2-Ti3Al1a-Ti) region has not been shown in Fig. 2(a).
As shown in Fig. 2(b), the selection area diffraction pattern (SADP) of TiN was indexed as a cubic unit cell with lattice pa-rameters of a 5 0.426 nm. The TEM/EDS analyses revealed that the layers of t2-Ti2AlN and t1-Ti3AlN were ternary compounds
with the approximate compositions Ti:Al:N 5 2:1:1 and Ti:Al:N 5 3:1:1, respectively. The first extensive study of phase equilibria in the Ti–Al–N system was conducted by Schuster and Bauer,30with two isothermal sections at 10001 and 13001C being
Fig. 4. After annealing at 13001C for 3 h: (a) bright-field (BF) image of the two-phase (a2-Ti3Al1a-Ti) region in Fig. 1(a); (b) dark-field (DF) image of a2
-Ti3Al formed by theð1120ÞTi
3Aldiffraction spot; (c) superimposed selected area diffraction patterns obtained from the two-phase (a2-Ti3Al1a-Ti) region,
showing the orientation relationship of½0001aTi j j½0001Ti3Alandð1100ÞaTi
j jð1100Þ
published. Recently, Pietzka and Schuster31showed that three ternary nitride phases Ti3Al2N2, Ti2AlN1x, and Ti3AlN1xare
present in a much more detailed section at 13001C. The SADPs of t2-Ti2AlN and t1-Ti3AlN are shown in Figs. 2(c) and (d),
respectively. On indexing these SADPs, it was found that t2-Ti2AlN had a hexagonal crystal structure with the lattice
parameters a 5 0.299 nm and c 5 1.361 nm, while the t1-Ti3AlN
was indexed to be a cubic unit cell with the lattice parameter a 50.4284 nm.
The a2-Ti3Al layer, as shown in Figs. 1(a) and 2(a), displayed
two distinct features, whose BF images are displayed in Figs. 3(a) and (b), respectively. The a2-Ti3Al layer had two different
morphologies, i.e., equiaxed a2-Ti3Al and elongated a2-Ti3Al.
Both equiaxed a2-Ti3Al and elongated a2-Ti3Al had the same
crystal structure but with different compositions (the Al con-centrations ranging from 22 to 39 at.%). The equiaxed structure marked ‘‘A’’ in Fig. 3(a) was corresponding to the regions also marked as ‘‘A’’ in Figs. 1(a) and 2(a), respectively. In the same way, the elongated texture structure marked ‘‘B’’ in Fig. 3(b) was corresponding to the regions both marked as ‘‘B’’ in Figs. 1(a) and 2(a). The EDS quantitative analyses indicated that the equiaxed a2-Ti3Al in Fig. 3(a) and elongated a2-Ti3Al in
Fig. 3(b) were 66.90 at.% Ti, 25.07 at.% Al, 8.03 at.% N, and 62.21 at.% Ti, 23.42 at.% Al, and 14.37 at.% N, respectively.
The elongated a2-Ti3Al grains in Fig. 3(b) were well aligned
nearly perpendicular to the interface of AlN/Ti. As Al and N diffuse much faster than Ti, the Ti region will be subjected to a state of compression because of the interdiffusion. However, the fact that the equiaxed a2-Ti3Al became elongated along the
di-rection perpendicular to the interface is attributed to the mis-match in the thermal expansion coefficient between AlN and Ti. Based upon our TEM investigation, the equiaxed a2-Ti3Al
was not randomly oriented, because it displayed the same contrast when the specimen was tilted. The equiaxed a2-Ti3Al
in Fig. 3(a) was likely to be formed by recrystallization. Previous studies have reported that the a2-Ti3Al alloys could exhibit
ad-mirable superplasticity of elongation greater than 1200% at temperatures between 9601 and 10001C and strain rates between 104and 105s1.32,33It is believed that the large deformation of textured a2-Ti3Al is likely to trigger the recrystallization.
Figure 3(c) displays the variation in the grain sizes of a2-Ti3Al
along the direction from the AlN side (bottom) to the Ti side (top). The grain size varies from 150 to 450 nm, implying the different degrees of grain growth after recrystallization. When the a2-Ti3Al, with a deformation texture, was recrystallized,
the new equiaxed grains usually had a preferred orientation.34 However, the orientation of the equiaxed a2-Ti3Al was different
from that of the textured a2-Ti3Al. The SADP, as shown
Fig. 5. (a) Bright-field (BF) image of g-TiAl between the t2-Ti2AlN and the lamellar structure after annealing at 15001C/0.5 h; (b) BF image of the
lamellar two-phase (g-TiAl1a2-Ti3Al) structure; (c) superimposed selected area diffraction patterns (SADPs) of g-TiAl (Z 5 [011]) and (Z 5 Z¼ ½2110)
from the lamellar structure region, showing the orientation relationship of½011TiAl j j ½2110Ti3Al; (d) schematic illustration of the SADPs in (c).
in Fig. 3(d), revealed that the orientation relationship of the equiaxed and the elongated a2-Ti3Al should be as follows:
½0001equiaxed j j ½1100elongated and ð1010Þequiaxed j j ð1122Þelongated. The fact that the exquiaxed a2-Ti3Al had a recrystallization
tex-ture was because of the influence that the textex-ture of the a2-Ti3Al
had on the nucleation and/or growth of the new grains. Figures 4(a) and (b) show the bright and dark field images of the two-phase (a2-Ti3Al1a-Ti) layer near the Ti side,
respec-tively. In Fig. 4(b), it can be seen that a-Ti is distributed along the grain boundaries of a2-Ti3Al. The lattice parameters of a-Ti,
calculated from the SADPs in Fig. 4(c), were a 5 0.310 nm and c 50.441 nm, while those of a2-Ti3Al were a 5 0.605 nm and
c 50.487 nm. The a value of a2-Ti3Al was approximately twice
that of a-Ti. It was inferred that the a2-Ti3Al was precipitated in
the matrix of a-Ti. From the SADPs, where the superlattice diffraction spots of a2-Ti3Al were clearly observed, the a2-Ti3Al
precipitates were found to satisfy the following orientation relationship with respect to a-Ti: ½0001aTi j j½0001Ti
3Al and
ð1100ÞaTi j j ð1100ÞTi
3Al. The orientation relations are
schemat-ically demonstrated in Fig. 4(d).
According to the Ti–Al binary phase diagram,35a2-Ti3Al was
stable up to 12101C, being a nonstoichiometric compound with a
relatively wide range extending from 23 to 35 at.%. It seemed that a2-Ti3Al could only be formed during isothermal annealing
below 12101C. However, the a2-Ti3Al phase was formed after
isothermal annealing at temperatures higher than 13001C, as mentioned previously. This implied that the a2-Ti3Al precipitated
from a-Ti during cooling. The existence of (a-Ti1a2-Ti3Al),
a2-Ti3Al, or (a2-Ti3Al1g-TiAl) various aluminide layers in the
reaction zone of the AlN–Ti diffusion couple, which was isother-mally annealed above 13001C, can be explained by a modified Ti–Al binary phase diagram, and will be described afterwards.
Figure 5(a) shows the BF image of the AlN/Ti interfacial re-action zone after annealing at 15001C for 0.5 h, showing t2
-Ti2AlN, g-TiAl, and the two-phase (g-TiAl1a2-Ti3Al) layer.
The inset on the upper right corner shows the SADP of g-TiAl with the zone axis of [110]. The (g-TiAl1a2-Ti3Al) lamellar
structure is displayed at a higher magnification in Fig. 5(b). The two-phase (g-TiAl1a2-Ti3Al) region usually forms a lamellar
morphology consisting of colonies of thin parallel a2-Ti3Al
and g-TiAl platelets. Based upon the Ti–Al phase diagram, the lamellar structure resulted from the eutectoid reaction of a-a21g. From the SADPs of the lamellar structure (Fig. 5(c))
and its corresponding schematic diagram (Fig. 5(d)), the
Fig. 6. (a) Isothermal section of the Ti–Al–N phase diagram at 13001C.30The diffusion path was drawn as arrows; (b) the microstructure of an AlN/Ti diffusion couple at 13001C; (c) the microstructure of an AlN/Ti diffusion couple after cooling.
orientation relationship of g-TiAl and a2-Ti3Al was identified:
½011TiAl j j ½2110Ti
3Alandð111ÞTiAl
j jð0110Þ
Ti3Al. In addition, the
schematic SADP patterns in Fig. 5(d) show some extra spots caused by the structure of twinning in tetragonal crystals of g-TiAl. The matrix of y-TiAl was oriented with its [011] zone axis parallel to the electron beam, while the twinning planeð111Þ, was in the edge-on direction along the electron beam. The dif-fraction patterns of the matrix and the twin were related by a mirror reflection across theð111Þ twinning plane or by a rota-tion of 1801 around the normal to the twinning plane.
(3) Microstructural Development and Diffusion Path at 13001C
When Ti comes into contact with AlN, the system becomes un-stable. It is generally acknowledged that the interfacial reactions between Ti and AlN include the following steps. AlN is reduced under the effect of Ti at the annealing temperature. The decom-posed Al and N atoms then diffuse into the Ti, and they react with each other to produce binary or ternary nitrides and alu-minides. While the nitrides are stable during the subsequent cooling, the aluminides will be subjected to phase transforma-tion based on the Ti–Al binary phase diagram.
Figure 6 illustrates the relationship between the Ti–Al–N phase diagram30 and the microstructure developed in the Ti/
AlN diffusion couple, which was isothermally annealed at 13001C. Figure 6(a) illustrates the isothermal section of the Ti–Al–N ternary phase diagram at 13001C. Based on the exper-imental results, it is proposed that the diffusion path, indicative of the compositions along the longitudinal direction perpendi-cular to the interface, was A–B–C–D–E–F–G–H when the dif-fusion couple was annealed at 13001C.
It is worth noting that the diffusion path is deviated from the ideal direct path between the compositions of the end points, i.e., Ti and AlN, as the diffusion velocities of the Ti, Al, and N atoms are significantly different in the Ti–AlN diffusion couple. The radii of N, Ti, and Al atoms are 0.07, 0.1448, and 0.1431
nm, respectively, so that N atoms, diffusing interstitially in the Ti, should have a higher diffusion velocity than that of Al at-oms, which diffuse mainly by substitution. In the literature, the diffusion coefficients of Al and N atoms in Ti are reported to be 7.4 107cm2/s at 6001–8501C and 1.2 102cm2/s at 9001–
15701C, respectively.36,37 Thus the TiN layer will firstly be formed at the interface, leading to the deviation of the diffu-sion path toward the TiN end (edge BC in Fig. 6(a)). As TiN has a wide range of nitrogen solubility and hardly dissolves Al at-oms, much more Al atoms than N atoms will diffuse into the titanium side beyond the TiN layer, resulting in the formation of t2-Ti2AlN, t1-Ti3AlN, and various aluminide layers.
The diffusion path crosses the fields of AlN1TiN, TiN, TiN1t2-Ti2AlN, t2-Ti2AlN1t1-Ti3AlN, t1-Ti3AlN1a-Ti, a-Ti,
a-Ti1b-Ti, and b-Ti. The tie lines of the ternary phase diagram correspond to the interface of two reaction layers in the diffusion couple, similar to the case for the interface between carbon and titanium aluminides presented by Viala et al.38 As shown in Fig. 6(b), the layers of TiN, t2-Ti2AlN, t1-Ti3AlN, a-Ti(Al,N),
a-Ti(Al,N)1b-Ti(Al,N), and b-Ti (Al,N) will be formed in se-quence from AlN to Ti at the annealing temperature (13001C). Figure 6(c) shows the microstructure formed during the cool-ing stage. It is believed that the nitride layers, includcool-ing TiN, t2
-Ti2AlN, and t1-Ti3AlN, remained during cooling. However, the
aluminide layers (i.e., a-Ti(Al,N), a-Ti(Al,N)1b-Ti(Al,N), and b-Ti(Al,N)) were subjected to phase transformation during the subsequent cooling, causing the formation of the a2-Ti3Al layer
and the lamellar two-phase (a2-Ti3Al1a-Ti) layer.
The phase transformation of aluminides, mentioned above, can be shown schematically in the Ti–Al binary phase diagram as shown in Fig. 7. Because nitrogen is an a-Ti stabilizers, the b/ (a1b) and (a1b)/a boundaries of the Ti–Al binary phase dia-gram are shifted upward by dissolving N atoms. The b-Ti layer abutting the reaction-unaffected Ti, with a small amount of Al and N in solid solution, was transformed into a-Ti on cooling, as indicated by line 1 in Fig. 7. The a-Ti1b-Ti layer at 13001C was transformed into the a-Ti1a2-Ti3Al layer after cooling (line
Fig. 7. A modified Ti–Al phase diagram because of the stabilization of a-Ti by dissolving N (see the dashed line),35showing the cooling processes of the various aluminides after annealing at 13001C (lines 1–3 ) and at 14001C (lines 4 and 5). The increase in the a-b transformation temperature has been exaggerated for clarification.
2 in Fig. 7), while the a-Ti layer was transformed into a2-Ti3Al
(line 3 in Fig. 7).
For comparison, a a-Ti layer and/or a g-TiAl layer with a higher Al content were likely formed when the diffusion couple was isothermally annealed at temperatures higher than 13001C. Thus, more aluminide layers, including the two-phase (g-TiAl1a2-Ti3Al) layer and/or g-TiAl layer, could exist in the
Ti/AlN diffusion couple because of the more extensive Al dif-fusion after annealing at 14001 or 15001C. After isothermal an-nealing at 14001C, the a-Ti layer with a high Al content underwent the following reaction: a-Ti-g-TiAl1a2-Ti3Al as
indicated by line 4 in Fig. 7, while the g-TiAl layer still remained as it was on cooling (line 5 in Fig. 7). It is worth noting that no g-TiAl existed in the diffusion couple after annealing at 15001C. After TiN1xwas saturated in N at such a high temperature, the
excess nitrogen would diffuse from the TiN1xlayer, resulting in
the nitridization of g-TiAl, that is t2-Ti2AlN.
(4) Reaction Zone Growth Mechanisms at 13001C
Based upon the results and the discussion mentioned above, an attempt was made to propose an interfacial reaction model be-tween AlN and Ti. The formation mechanisms for several
dif-ferent stages at 13001C, as an example, are schematically shown in Fig. 8.
(A) First Stage: Formation of TiN: In the first stage, AlN decomposes at the original AlN/Ti interface into Al and N under the effect of Ti. As Ti has a strong affinity with N, the fast diffusion species N atoms will go into Ti, leading to the forma-tion of TiN with Al in solid soluforma-tion. The formaforma-tion of TiN in the initial stage can be expressed by the following reactions, as shown in Fig. 8(a):39
AlN! Al þ N DG¼ 159; 770 ðJ=moleÞ (1)
Tiþ N ! TiN DG¼ 188; 108 ðJ=moleÞ (2)
AlNþ Ti ! TiN þ Al DG¼ 347; 878 ðJ=moleÞ (3) The diffusion of N atoms and Al atoms in TiN is much faster than that of Ti atoms in AlN (diffusivities of N and Al in TiN are 5.4 103 cm2
/s at 10001–15001C and 3 1014 cm2
/s at 3001–5501C; the diffusivity of Ti in AlN is 4 1017at 12801–
14001C).40 Thus, Al and N continuously diffuse through the TiN layer and reaction (2) takes place at the interface of TiN and Ti. The highly negative Gibbs free energy change of the reaction (3) indicates that the reaction between aluminum nitride and titanium is thermodynamically favorable.
(B) Second Stage: Formation of Ti3AlN and Various Titanium Aluminides: When the concentrations of Al and N increase to a certain amount, t1-Ti3AlN is formed, as shown in
Fig. 8(b), at the interface of TiN/Ti according to the following reaction:
3Tiþ Al þ N ! t1-Ti3AlN (4)
The formation of t1-Ti3AlN instead of t2-Ti2AlN at this stage
is consistent with the fact that t1-Ti3AlN was found after
an-nealing at 13001C/0.5 h, which was at a relatively early stage. Meanwhile, excess Al and N atoms further go into Ti as a solid solution, leading to the formation of various titanium alumi-nonitrides, for example, a-Ti(Al,N) and the a-Ti(Al,N)1 b-Ti(Al,N) two-phase region.
(C) Third Stage: Formation of Ti2AlN: As shown in Fig. 8(c), t2-Ti2AlN is formed at this stage. The growth of the
t2-Ti2AlN was controlled by the diffusion of Al and N through
TiN, accumulating at the TiN/t1-Ti3AlN interface, leading
to the disappearance of t1-Ti3AlN based upon the following
reaction:
t1-Ti3AlNþ12Alþ21N!32t2-Ti2AlN (5)
(D) Fourth Stage: Formation ofa2-Ti3Al and/or the Two-Phase Layer (a2-Ti3Al1a-Ti) During Cooling: Figure 8(d) shows that the a-Ti(Al,N)1b-Ti(Al,N) two-phase region was transformed into the (a2-Ti3Al1a-Ti) two-phase layer, while the
a-Ti(Al,N) layer was transformed into the a2-Ti3Al layer during
cooling, depending upon the local composition. Not shown in Fig. 8(d) is that the b-Ti(Al,N) was transformed into the a-Ti(Al,N). The phase transformations mentioned above can be expressed by the following reactions,
b-TiðAl; NÞ !coolinga-TiðAl; NÞ (6)
a-TiðAl; NÞþb-TiðAl; NÞ !coolinga-TiðAl; NÞþa2-Ti3AlðNÞ (7)
a-TiðAl; NÞ !coolinga2-Ti3AlðNÞ (8)
Ti
(c) AlN Ti Ti3AlN Ti2AlN β-Ti(Al, N) TiN two-phase region AlN TiN Ti Ti3AlN Ti2AlN Ti3Al (d) AlN Ti Al TiN N (a) Ti AlN Ti Al TiN Ti3AlN (b) N β-Ti(Al, N)Fig. 8. Proposed reaction mechanisms of the AlN and Ti diffusion cou-ple annealed at 13001C: (a) first stage: formation of TiN; (b) second stage: formation of t1-Ti3AlN, a-Ti(Al,N)1b-Ti(Al,N), and a-Ti(Al,N);
(c) third stage: formation of t2-Ti2AlN; (d) fourth stage: formation of
a2-Ti3Al and/or the two-phase layer (a2-Ti3Al1a-Ti) during cooling.
IV. Conclusions
1. An interfacial reaction zone, consisting of TiN, t2
-Ti2AlN, t1-Ti3AlN, a2-Ti3Al, and a two-phase (a2-Ti3
Al1a-Ti) region in sequence, was observed in between AlN and Ti after annealing at 13001C.
2. After annealing at 13001C, a textured structure existed in the a2-Ti3Al layer probably because of the internal stresses
re-sulting from the mismatch in the thermal expansion coefficient between AlN and Ti. The fine equiaxed a2-Ti3Al grains implied
the occurrence of recrystallization. The orientation relationship between the equiaxed and elongated a2-Ti3Al was as follows:
½0001equiaxed j j ½1100elongatedandð1010Þequiaxed j j ð1122Þelongated. 3. In the two-phase (t1-Ti3AlN1a-Ti) region after
anneal-ing at 13001C, a2-Ti3Al and a-Ti were found to satisfy the
fol-lowing orientation relationship: ½0001aTi j j½0001Ti
3Al and
ð1100ÞaTi j jð1100ÞTi3Al. The a value of a2-Ti3Al was
approxi-mately twice that of a-Ti.
4. The g-TiAl and a lamellar two-phase (g-TiAl1a2-Ti3Al)
structure were found in between t2-Ti2AlN and a2-Ti3Al after
annealing at 14001C. The orientation relationship of g-TiAl and a2-Ti3Al in the lamellar structure was identified to be as follows:
½2423T
3Al
j j½012 j j
TiAl and ð1010ÞTi3Al j j ð100ÞTiAl. Compared
with the results after the reaction at 14001C, g-TiAl was not formed at the interface after the reaction at 15001C.
5. The diffusion path, connecting the phases formed by the reaction between AlN and Ti, was drawn on the Ti–Zr–O ternary phase diagram. Furthermore, the relationships among the Ti–Al–N ternary phase diagram, a modified binary phase diagram, and the microstructural development between AlN and Ti at 13001C and subsequent cooling have been pro-posed.
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