Processing and high temperature deformation of Al2O3 fiber-reinforced NiAlFe matrix composites

ELSEVIER

Materials Science and Engineering

A191 (1995) 49-59

Processing and high temperature deformation of A&O, fiber-reinforced NiAlF’e matrix composites Cyrus Liua, S.M. Jeng”, J.-M. Yang”, R.A. Amatob aDepartment of Materials Science and Engineering, University of California, Los Angeles, CA 90024-1595, b General Electric Aircraft Engines, Cincinnati, OH 4.5125, USA

USA

Received 7 March 1994; in revised form 20 June 1994

Abstract The elevated temperature mechanical behavior and deformation composites fabricated by powder spray and tape casting techniques temperature, 760, 870, and 980 “C. Flexural creep tests were conducted role of fiber, matrix, and the interface on the tensile and creep behavior of Keywords: Deformation;

Aluminium;

characteristics of A&O, fiber-reinforced NiAlFe matrix were investigated. Tensile tests were conducted at room at 700 “C with stresses ranging from 100 to 250 MPa. The the composite is also discussed.

Oxygen; Nickel; Iron; Composites

1. Introduction Ordered intermetallic compounds have emerged as a new class of materials for advanced structural applications. NiAl-based intermetallics, in particular, have been recognized as one of the most promising candidate materials for high temperature applications. They possess several attractive properties including low density (about 6 g cm- 3), high melting point ( 1638 “C), high modulus ( 189 GPa) and excellent oxidation resistance up to 1300 “C. Polycrystalline NiAl exhibits a brittle-to-ductile transition at temperatures ranging from 300 to 600 “C, the exact temperature depending on the stoichiometry and grain size. However, to make NiAl a viable material, it is necessary to overcome some of its inherent problems. These include low ductility and fracture toughness at ambient temperatures and inadequate strength and creep resistance at elevated temperatures. Accordingly, significant efforts have centered on enhancing the mechanical properties of NiAl through grain refinement and micro and macro alloying as well as incorporating second-phase reinforcements [l-4]. Recent studies have focused on using a composite approach in order to improve the poor high temperature properties of NiAl. Single-crystal aluminium oxide 0921.5093/95/$9.50 0 1995 - Elsevier Science S.A. All rights reserved SD1 0921-5093(94)09634-9

fibers have been shown to be one of the most promising reinforcements for several intermetallic matrix composites [5]. In addition to good chemical compatibility with NiAl, these fibers also possess excellent high temperature creep resistance [6]. Based on the properties of the constituents, the resulting A&O,-NiAl composite should have improved high temperature properties when compared with monolithic NiAl. The degree of improvement, however, will depend on the strength of the inter-facial bond at elevated temperatures. While a strong inter-facial bond may improve load transfer from the creeping matrix to the more creep-resistant fibers, others have shown that the use of weak interfaces offer increased toughness in addition to improved creep resistance [ 71. This work was conducted to study the high temperature deformation of an A&O, fiber-reinforced NiAl matrix composite. The effect of different processing conditions on the elevated temperature properties of these composites was also investigated.

2. Experimental procedure The matrix alloy used was a single-phase (p) alloy, consisting of 32 at.% Al, 20 at.% Fe, and the balance

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Ni. Macroalloying NiAl with Fe has been shown to increase room temperature ductility compared with stoichiometric NiAl [8,9] in addition to decreasing the steady state creep rate [lo]. Single-crystal Al,O, fibers produced by SaphikonrM (Milford, NH) were used as reinforcing materials. These fibers have an average diameter of 140 pm and a tensile strength of 3.15 GPa at room temperature. Six-ply unidirectional A&O,NiAlFe composites were fabricated using two different powder metallurgy routes. In the first method, individual plies with approximately 50 fibers in-’ were produced by spraying NiAlFe powder mixed with a polymeric binder onto alumina fibers which were placed in a 6 in X 6 in grooved plate. To fabricate these powder spray composites, the required number of layers were laid up and then consolidated using hot isostatic pressing (HIP) at 1120 “C (2050 OF) and 172 MPa (25 klbf in’) for 2 h. The use of a two-stage HIP cycle permits binder removal. The other technique used to fabricate the composites used in this study involved the use of tape casting [ 111. In this process, the matrix is again mixed with a polymeric binder. The resulting slurry is then spread over a fiber mat and cast to form a composite monotape of desired thickness and size. The desired number of layers are then laid up, degassed, and subjected to, in this case, vacuum hot pressing (VHP) at a temperature of 1120 “C (2050 F) using a pressure of 24 MPa (3.5 klbf in2) for 2 h and HIP at 1120 “C (2050 “F) and 172 MPa (25 klbf in-‘) for 2 h to achieve final consolidation. A two-stage VHP cycle is used first to burn off the binder and second to achieve consolidation. The HIP step completes composite densification. A sample of monolithic NiAlFe alloy was also fabricated using tape casting and consolidated under the same conditions. In order to examine their microstructures, samples were cut from the composite materials, prepared according to standard metallographic procedure, and examined using an optical microscope. Thin sections of varying thicknesses (432-533 pm) were also cut from both the powder spray and tape cast composites, using a low speed diamond saw, and polished to a 0.05 pm Al,O, finish. These specimens were used to measure the interfacial shear sterngth of the composite via an indentation (pushout) test [ 121. In this test, a microhardness tester is used to push out the individual fibers in order to estimate the strength of the interfacial bond. An average of 22 fibers were pushed out of each of the specimens. A scanning electron microscope was used to examine the morphology of the debonded interface. Tensile tests were conducted at room temperature, 760 “C (1400 OF), 870 “C (1600 F), and 980 “C (1800 F) using an Instron servohydraulic testing machine operated at a cross-head speed of 0.5 mm min’ on specimens 50 mm long by 12.5 mm wide by

3.15 mm thick with the fibers parallel to the longitudinal direction; the dimensions of the gauge section were 25.4 mmX 6.25 mm. The strain to failure was monitored using an extensometer. A minimum of three specimens were tested at each temperature. Four-point bending creep tests were performed at 700 “C in a vertical split-tube furnace on pure NiAlFe and A&O,-NiAlFe composites with the approximate dimensions of 50 mm X 5 mm X 5 mm. The deflection at the center of the beam was measured with a threeprobe linear variable-deflection transducer. The transducer output was routed through a digital multimeter to a microcomputer which converted the signal to give the strain in the sample in real time. The experimental set-up and data reduction method can be found in Ref.

1131. Fractographic analysis was performed on all tested specimens using a scanning electron microscope. Specimens which failed in tension were further polished along their length to reveal the extent of damage in the areas near the fracture surface. This allowed for determination of the mechanisms of damage initiation and propagation.

3. Results and discussion 3.1. Microstructure and intellfacial mechanical properties Figs. l(a) and l(b) show the microstructure of the powder spray Al,O,-NiAlFe composite along the transverse and longitudinal directions respectively. The composites were fully dense with the Al,O, fibers quite well aligned in the matrix. The grain boundaries can be clearly seen in the micrograph together with the variation in the grain size of the matrix material. The fiber volume fraction was estimated to be approximately 10%. Fiber fractures, evident in Fig. l(b), are believed to have occurred during consolidation of the composite. The high HIP pressure could have broken the fibers and permitted matrix particles to infiltrate between the fiber ends. Fig. 2 is a scanning electron microscopy micrograph of the inter-facial region in the powder spray A&O,-NiAlFe composite. No evidence of reaction between the fiber and matrix was observed in the composites; however, some voids were found at the interface. The microstructure of a tape cast Al,O,-NiAlFe composite along the transverse direction is shown in Fig. 3(a). Similar to the powder spray composites, these composites seemed to be dense with some voids within the grains (see Fig. 3(b)) and exhibiting a range of grain sizes. On the contrary, the fibers exhibited much better alignment and no evidence of fiber damage resulting from consolidation was found.

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Material Science and EngineeringAl

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Fig. 2. Microstructure of the interfacial region in a powder spray A&O,-NiAlFe composite.

1200 pm Fig. 1. Microstructure of a powder spray Al,O,-NiAlFe composite along the (a) transverse and (b) longitudinal directions.

The volume fraction of fibers varied beteen 15.7% and 20.7%. The room temperature interfacial shear strengths of the A&O,-NiAlFe composites fabricated by the two different techniques are listed in Table 1. The results indicate that the tape cast material exhibited a higher inter-facial shear strength (240 MPa) than the powder spray composite (136 MPa). Several previous studies [ 14-161 also indicated that the processing routes affected the inter-facial bond strength in Al,O,-NiAl composites, as shown in Table 1. Figs. 4(a) and 4(b) show the inter-facial debondmg locations of the powder spray and the tape cast A&O,-NiAlFe composites respectively. In the powder spray A&O,-NiAlFe composite, the fiber surface is quite clean with only a small amount of matrix material adhering. The fibers pushed out from the tape cast material (Fig. 4(b)) appeared very similar to those from the powder spray

Fig. 3. Microstructure of a tape cast A120,-NiAlFe composite: (a) transverse direction; (b) grain structure, showing voids within the grains.

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Table 1 Interfacial shear strengths

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Material Science and Engineering Al 91 (1995) 49-59

r, of Al,O,-NiAl-based

composites

Material

Al,%NiAlFe”

Al,%NiAIFeb

Al,%NiAl”

r, (Mf’a)

136k8 -

240f19 -

25-100

50-150

> 280

53-248

62-182

[I31

[I41

[I41

1151

[I51

Ref.

Al@NiAl’

Al@,NiAld

Al,% NiAL’

Al,% NiAI’

Al@,NiAlg +

158-190+

[I51

BProcessed using the powder spray technique and HIP. h Processed using tape casting, VHP, and HIP. c Processed using the powder cloth process, hot pressing, and HIP. d Processed using hot pressing and HIP without the use of binders. r Processed using casting. ’Processed using zone directional solidification. S Processed using zone directional; thermally cycled after processing.

will not occur until 1310°C (1583K) [17], it is unknown whether this result would apply to NiAlFe. In addition, no evidence of chemical attack was observed on the fiber surface. Although energy-dispersive X-ray spectroscopic analysis of the material adhering to the fiber surface did not yield any results, a detailed transmission electron microscopy examination of the interfacial region may provide further insight. In addition, there might be a large frictional contribution to ri, probably because of residual clamping and a mechanical interlocking effect resulting from post-consolidation cooling and the large difference in coefficients of thermal expansion (CTES) between the matrix and fiber (15X 10m6m m-l K-l [18] and 9~ 10P6m m-l K- * [ 191 respectively). It is expected that the inter-facial shear strength and frictional stress would be significantly reduced at elevated temperatures owing to a reduction in the thermomechanical clamping forces ]201. 3.2. Tensile behavior

Fig. 4. Interfacial debonding location in (a) powder spray and (b) tape cast A&O,-NiAlFe composites.

composites, exhibiting some adhering material on their surface. The material adhering to the surface of the fiber seems to suggest that some chemical reaction has occurred between the NiAlFe matrix and the A&O, fibers. While thermodynamic calculations have indicated that chemical reactions between A&O, and NiAl

The tensile properties of monolithic NiAlFe, powder spray, and tape cast Al,O,-NiAlFe composites tested at 25, 760,870, and 980 “C are listed in Table 2. At room temperature, both the monolithic NiAlFe alloy and the A&O,-NiAlFe composites exhibited linear elastic stress-strain behavior until fracture. However, when tested at elevated temperatures, NiAlFe and the composites exhibited extensive plastic deformation before fracture. Similar stress-strain curves were also measured for A&O, fiber-reinforced NiAl composites fabricated by the powder cloth approach [ 151. The strain to failure of the NiAlFe alloy increased from less than 1% at room temperature to approximately 20% at 870 “C. The room temperature ultimate tensile strength of the monolithic alloy is higher than that of both composites. This may be due to the loss of fiber strength caused by the high proces-

C. Liu et al. Table 2 Tensile properties Material

/

of NiAlFe and the AI@-NiAlFe

composites

Elastic modulus (Gpa)

Ultimate tensile strength (MPa)

Room temperature

760 “C

870 “C

NiAlFe A&O,-NiAIFe”

136 152

88 102

68 77

AI@-NiAlFe”

193

-

-

.1Processed using the powder spray technique ” Processed via tape casting, VHP, and HIP.

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Material Science and Engineering A191 (1995) 49-59

980 “C

96

Room temperature

760 “C

870 “C

980 “C

422 333

160 181

85 96

-

405

-

-

96

and HIP.

sing temperatures used during consolidation of the composites and the strong interfacial bond strength in the composites which led to behavior similar to that in fiber-reinforced ceramic matrix composites with strong interfaces. In addition, the powder spray composites suffer from fiber breakage during consolidation, reducing the effective length of the A&O, fibers. However, at elevated temperatures, the ultimate tensile strengths and elastic moduli of both composites are higher than that of the pure matrix alloy. The fracture surfaces of monolithic NiAlFe tested at three different temperatures (25 “C, 760 “C, and 870 “C) are shown in Figs. 5(a)-(c) respectively. They clearly show that the fracture mode of the NiAlFe alloy has transformed from transgranular cleavage at room temperature to intergranular fracture at intermediate temperatures and, finally, to ductile dimple failure at high temperatures. Figs. 6(a)-(c) show the tensile fracture surfaces of powder spray Al,O,-NiAlFe composites after testing at the same three temperatures. They clearly show that the fiber pullout length of the composites increased as the testing temperature increased. The strong interfacial shear strength of the composite at room temperature prevents any fiber pullout, resulting in catastrophic failure with virtually no fiber pullout. However, as the testing temperature increased, the residual clamping stress resulting from the CTE mismatch will decrease, leading to a reduction in the interfacial shear strength and allowing debonding and fiber pullout from the matrix to occur as the loads increase. Meanwhile, transgranular cleavage and intergranular failure were observed in the matrix near the inter-facial region of the composites tested at 760 “C and 870 “C respectively (Fig. 7). These failure modes were quite different from those of the monolithic alloy tested at the same temperatures, Furthermore, examination of the microstructure beneath the fracture surface sowed that multiple fiber fractures were observed only in the composite tested at high temperature. Meanwhile, matrix cracking (Fig. S(a)) and relative displacement between the fiber and matrix are also found near the

broken fiber end (Fig. 8). The fracture morphology of the tape cast composites was similar to that of the powder spray material. 3.3. Creep behavior Fig. 9(a) shows the creep curves of the monolithic matrix and powder spray A&O,-NiAlFe composite tested at 700 “C. Both materials clearly exhibit primary and steady state creep behavior. For the pure matrix material, the steady state creep rates at 72 MPa, 100 MPa, and 136 MPa are 7 X lo-* s-l, 2.64 x lo-’ SK’, and 4.1 x lo-’ s-’ respectively. The secondary creep rates in the powder spray composite at 100 MPa and 136 MPa are 1.86~ lo-’ SK’ and 3.66~ 10 s-i respectively. The fact that the powder spray composite behaved in a similar fashion to the monolithic NiAl alloy can be explained by the fact that broken fibers do not appreciably change the mechanical properties of the unreinforced matrix alloy. Fig. 9(b) shows the typical creep curves of tape cast A&O,-NiAlFe composites tested at 700 “C. The behavior of the composites exhibits the three classic creep stages. Initially, large deformation occurs in the primary creep region, followed by a long secondary creep region. In general, very little deformation was observed in this regime; in fact, the initial results seemed to imply that inverse creep was occurring. For applied stresses of 150 MPa and 175 MPa, the tape cast material had steady state creep rates of 5.12 x 10-l” s-’ and 6.07 x 10-l” s-i respectively, a reduction of three orders of magnitude compared to the powder spray composite. It is obvious that the creep resistance of NiAlFe can be improved through the incorporation of Al,O, fibers. The large difference in steady state creep rate between the two composites is due to the stronger interfacial bond and the higher fiber volume fraction in the tape cast material. The stronger interfacial bond will improve load transfer from the matrix to the fibers while the higher volume fraction of A&O, fibers markedly improves the creep strength of the material. In addition, the unbro-

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Fig. 5. Tensile fracture surfaces of monolithic (a) 25 “C,(b) 760 “C, and (c) 870 “C.

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Material Science and Engineering A191 (1995) 49-59

NiAlFe tested at

ken fibers in the tape cast composites are more effective reinforcements than the broken fibers in the powder spray composites. At applied stresses higher than 175 MPa, however, the steady state creep rate appears to increase rapidly, as shown in Fig. 10. This

Fig. 6. Tensile fracture surfaces of powder-spray Al,O,-NiAlFe composites tested at (a) 25 “C,(b) 760 “C, and (c) 870 “C.

sensitivity of the steady state creep rate to the applied stress was also observed in the creep of Sic fiber-reinforced Ti matrix composites [13]. Failure, which occurred in composites with applied stresses above 175 MPa, followed a short period of tertiary creep in

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Material Science and Engineering Al 91 (1995) 49-59

Fig. 7. Failure morphology near the interfacial region in powder spray A120,-NiAIFe composites tested at (a) 760 “C and (b) 870 “C.

which a large amount of deformation took place in a very short period of time. Examination of tape cast composites which failed under creep revealed a change in the deformation mechanisms for stresses greater than 175 MPa. When the applied stresses were lower than 175 MPa, most of the Al,O, fibers remained intact, as shown in Fig. 1 l(a), leading to a final strain of approximately 0.55%. Cavitation within the grains of the matrix material was observed together with the presence of many voids at the grain boundaries, as seen in Fig. 12. Grain boundary diffusion is thought to be responsible for enlarging these voids. Because the alumina fibers remained intact in this stress regime, creep of the composite was limited despite relaxation of the matrix. When a,,,, was greater than 175 MPa, however, multiple fiber breakage occurred during the secondary creep stage (E = 0.55% for a,,,, = 200 MPa; E = 0.61% for c~+,~,= 253 MPa), as seen in Figs. 1 l(b) and 1 l(c). Increasing the applied stress increased the number of broken fibers (see Fig.

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Fig. 8. Microstructure near the broken fiber end in powder spray ALO,-NiAIFe composites tested at (a) 760 “C and(b) 870 “C.

1 l(c)). The fiber failures led to matrix cracking, resulting in the eventual failure of the composite. The fracture surfaces of a crept tape cast Al,O,-NiAlFe composite, shown in Fig. 13, exhibit the typical failure modes of brittle materials. Intergranular fracture in the matrix was combined with cleavage failure of the Al,O, fibers. Exposure to air during the tests have formed oxidation products on the grain boundaries (see Fig. 13(b)). There was also no noticeable pullout of the fiber, as seen in Figs. 13(a) and 13(c), indicating the interfacial bonding in this composite remained quite strong. This was also observed when the fiber-matrix interfaces in the tape cast composites were examined after creep testing (Fig. 14). Despite the fact that the strength of the interfacial bond should have decreased owing to extended exposure at high temperatures, there seems to be no evidence of debonding in the composite tested at 200 MPa (Fig. 14(a)).

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o.ooof 0

’ 3000

c

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Material Science and Engineering Al 91 (1995) 49-59

6000

9000

12000

15000

16000

Time. I (minutes) (b)

Fig. 9. Creep curves of (a) monolithic NiAlFe and powder spray A120,-NiAlFe composites and (b) tape cast A&O,-NiAIFe composites tested at 700 “C.

The difference in creep behavior for applied stresses above 175 MPa, therefore, is believed to be due to fiber breakage. The strength of the as-received fiber has been shown to decrease from 3 GPa at room temperature to about 0.8 GPa at 1120 “C [6], the temperature used for composite consolidation, On being subjected to creep loading, the fiber is believed to carry the majority of the load compared with the matrix. However, as the matrix gradually relaxes during creep deformation at 700 “C, the stress in the fiber would increase. The high stresses in the fibers would lead to the occurrence of random fiber fracture due to the statistical nature of fiber strength. At the fiber rupture sites, the elastic load released by the fiber is transferred back to the matrix which sustains further creep deformation. As the matrix continues to creep, the local stress is continuously redistributed until an intact fiber adjacent to or a fiber segment further along the ruptured fiber supplies a constraining force allowing the stress in the matrix to relax again. As the process continues, multiple fiber failure is the resulting microstructure, as illustrated in Figs. 12(b) and 12(c). This build-up of stress in the matrix due to fiber fracture will lead to the initiation of matrix cracking. Crack development resulting from the linkage of matrix cracking and fiber failure would lead to the final failure of the composite.

4. Summary

106

I 0

tape-cat

Al,OfiiAlFe

I

IO2 Applied Stress, o (MPa) Fig. 10. The relationship between the steady state creep rate and applied stress for monolithic NiAlFe and powder spray and tape cast AllO,-NiAIFe composites tested at 700 “C.

Two different processing routes, powder spraying and tape casting, were used to produce unidirectional A&O,-reinforced NiAlFe composites. The former method involved spraying NiAlFe powder mixed with a binder onto an A120, cloth to form a composite ply which was then stacked and consolidated. The latter technique deposited a binder-powder slurry into fiber cloths to form monotapes which were laid up and consolidated to obtain a composite. Tape casting appears to be a much more promising technique for fabricating intermetallic matrix composites. In addition to producing monolayer tapes with a higher volume fraction of fibers than the powder spray method, tape casting also produces more flexible tapes, enabling more complex structural components to be fabricated. Another benefit of tape casting is better fiber alignment after consolidation. The Al,O,-NiAlFe composite produced using the tape casting process showed a higher interfacial bond strength than that obtained using the powder spray technique. This, coupled with the higher volume fraction, is responsible for the substantially improved creep behavior or the tape cast material over the powder spray composite. There is, however, still no

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Material Science and Engineering Al 91 (1995) 49-59

Fig. 11. Microstructure of tape cast A&O,-NiAlFe *;,pp1= 200 MPa; (c) a,,,, = 253 MPa.

composites

after four-point

57

creep testing at 700 “C: (a) oappl= 150 MPa; (b)

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12. Microstructure of the matrix in a tape cast A&O,-NiAlFe composite loids collect at the grain boundaries; (b) cavitation within a grain.

13. The fracture surface of a crept tape cast A120,-NiAlFe wing oxidation products; (c) Al,O, fibers.

composite

after four-point

creep testing (a,,,, = 150 MPa) at 700 “C:

( T= 700 “C, a,,,, = 150 MPa): (a) overall view; (b) matrix,

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tion, the elevated temperature fibers need improvement.

59

strength of the Al,O,

Acknowledgment

This work was partially supported by the National Science Foundation (DDM 905 1030).

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1:;

Fig. 14. The fiber-matrix interface in a tape cast Al?O,-NiAlFe composite after creep testing at 700 “C: (a) qlpp, = 200 MPa; (b) u dpp,= 150 MPa.

concrete evidence to indicate that there is a chemical reaction contributing to the high interfacial bond strength of the tape cast material. For NiAl-based intermetallic matrix composites to become viable structural materials, the poor creep resistance of the matrix material must be overcome via alloying. In addi-