Ni3Al FGMs

Plh S 1359-8368(96)00027-3

ELSEVIER

Composites Part B 28B (1997) 21 27 ~7~1997 ElsevierScienceLimited Printed in Great Britain. All rights reserved 1359-8368/97/$17.00

Cyclic thermal shock resistance of TiC/Ni3AI FGMs

L. M. Zhang and T. Hirai Institute for Materials Research, Tohoku University, Katahira 2- 1- 1, Aoba, Sendai 980-77, Japan

and A. K u m a k a w a Kakuda Research Center, National Aerospace Laboratory, Kimigaya, Kakuda, Miyagi 981- 15, Japan

and R. Z. Yuan Institute for Advanced Materials Research, Wuhan University of Technology, Wuhan 430070, P. R. China (Received 21 December 1995; revised 22 March 1996) Cyclic thermal shock tests on three TiC/Ni3A1 FGM specimens were carried out under simulated largetemperature-difference conditions. The effective thermal conductivity was measured and the variation of microstructure was observed. The effects of the sample geometry and the TiC side surface state on the cyclic thermal shock behavior were investigated, together with thermal stresses under steady-heating conditions. The relation between the variation of effective thermal conductivity and crack propagation was also investigated. The damage which occurred in these FGMs was mainly caused by excessive inplane compressive stresses on the surface of the TiC side during the cyclic thermal shock tests. © 1997 Elsevier Science Limited. All rights reserved (Keywords: cyclic thermal shock; TiC/Ni3AI; functionally graded materials)

1 INTRODUCTION Thermal stress relaxation in ceramic/metal type F G M s (functionally graded materials) is believed to hold promise for applications in the thermal protection systems of space planes and in the innermost wall of thermonuclear experimental reactors 1, where the materials must endure high temperatures and cyclic heat loads under temperaturegradient conditions. Therefore, thermal shock characteristics of F G M s should be evaluated according to different research aims. Recently, evaluations of heat-resistant properties of oxide/metal and carbide/carbon FGMs have been carried out with a laser shock test system, a xenon arc-lamp/liquid nitrogen test system, and a gas burner test system, all of which simulate the practical circumstances found in the space plane 2m. TiC/Ni3A1 F G M specimens for potential use as the innermost wall materials of thermonuclear experimental

reactors have been successfully fabricated by the authors 5. The innermost wall is in direct contact with the plasma, so the materials must be of low atomic number and have high temperature resistance and excellent cyclic thermal shock resistance 6'7. In this paper, the cyclic thermal shock characteristics of TiC/Ni3A1 FGMs were tested under large temperature-difference conditions with a xenon arclamp/liquid nitrogen test system. The relationships among the variations in effective thermal conductivity, damage, and morphology of the F G M s were also investigated.

2 EXPERIMENTAL 2.1 Design and fabrication on TiC/Ni3Al FGMs

Three disk-shaped specimens, WU-80, 81 and 82, were designed and prepared according to the rule of minimum stress ratio in the residual stress of TiC/Ni3A1 FGMs.

21

Cyclic thermal shock of TiC/Ni3AI FGMs: L. M. Zhang et al. The specimens had the same diameter, D = 30 mm, the same constituent distribution exponents, P = 1.2 (P = In c/(ln x - in d), where c is the composition of Ni3A1, d is the total thickness of the graded interlayers, and x is the layer location coordinate measured from the TiC side); and the same thickness di = 0.5 mm of each interlayer in the graded region. The specimens differed in their bulk thickness and in their structure around the TiC side. The thicknesses of specimens WU-80, 81, 82 were 6.5 mm, 5.0 mm, and 6.0 mm, respectively. TiC, Ni3A1 and Cr3C 2 powders with average particle size of 5, 6.1 and 6.5 #m, respectively, were used in this study. Table 1 shows a summary of the test specimens, dimensions and composition. Both WU-80 and WU-82 specimens consisted of 11 layers. In the WU-80 specimen, 3 vol% of Cr3C2 as a sintering aid was added to the TiC layer and to the TiCrich layers containing less than 20vo1% of Ni3A1 to promote the densification of the ceramic side 8. The WU81 specimen was composed of 10 layers, each layer having a thickness of 0.5 mm; 3 vol% of Cr3C 2 was also added to the TiC layer and TiC-rich layers. No Cr3C2 was added to the WU-82 specimen, resulting in the formation of porous layers near the TiC side, in order to examine the effect of such layers on thermal shock. The mixtures were hot-pressed at 1573 K and 30 MPa for 2 h in gaseous argon atmosphere. Specimens were brazed with copper plates with a thickness of 2mm at 1073K for 15min, as shown in Figure 1. A commercially available solder containing Ag, Cu and Ti was used. The compositional distribution and layer structure of the WU-80 specimen are shown in Figure 2.

copper substrate which was cooled by liquid nitrogen so that the specimen could be steadily subjected to largetemperature-differences by heating the TiC surface with a beam of light from the xenon arc-lamp. Twenty heating-cooling cycles were achieved by opening and closing a shutter. The heat flux was maintained at about 1.5 MW/m 2.

2.3 Observation of microstructure after cyclic thermal shock tests The morphology and microstructure of the surface and vertical cross-section after thermal shock tests were observed under an optical microscope. The hardness distributions along the vertical center line and along a parallel line at a position of about 2/3 of the radius of the cross section were measured with a Vickers hardness instrument.

2.2 Apparatus of cyclic thermal shock tests Figure 3 is a schematic diagram of the test apparatus, consisting of a vacuum system, a sample-holder, a 30 KW-xenon arc-lamp for heating, two reflecting mirrors, a liquid nitrogen cooling system, and a test-control system. The Ni3A1 side of the specimen was brazed to the

Figure 1 Specimen of a TiC/Ni3AI FGM (P = 1.2)

Table 1 The constitution of TiC/Ni3A1 FGMs

Layer number

No. 1

No. 2

No. 3

No. 4

No. 5

No. 6

No. 7

No. 8

No. 9

No. 10

No. 11

0.5*

0.5*

0.5

0.5

0,5

0.5

0.5

0.5

1.0

26.8

35.0

43,5

52.4

61.5

80.4

0.5

0.5

0,5

0.5

0.5

0.5

WU-80 (Thickness: 6.5 ram) Thickness of layers (rnm) 1.0' Content of Ni3A1 (vol%) 0.0

0.5*

WU-81 (Thickness: 5.0mm) Thickness of layers (mm) 0.5* Content of Ni3A1 (vol%) 0.0

0.5* 5.1

11.7

19.0

26.8

35.0

52,4

61.5

80.4

WU-82 (Thickness: 6.0 mm) Thickness of layers (ram) 0.5 Content of Ni3AI (vol%) 0.0

0.5

0.5

0.5

0.5

0.5

0,5

0.5

0.5

0.5

5.1

11.7

19.0

26.8

35.0

43.5

52.4

61.5

80.4

* 3 vol% Cr3C2 was added

22

5.1

11.7

0.5*

19.0

0.5*

100

100

1.0 100

Cycfic thermal shock of TiC/Ni3AI FGMs: L. M. Zhang 2.4 Thermal stress analysis under steady-heating conditions The thermal stress distributions in FGMs at steady temperature difference conditions during the cyclic heatingcooling process were analyzed by means of the finiteelement method. The calculation model and mesh geometry used in this study were similar to those used for the second optimization design of Ni3A1/TiC FGMs 9, except that the brazed Cu plate was as shown in Figure 1. The side of the Cu plate was restrained a~ad insulated heat conditions at the free boundary of the specimen were assumed. The measured maximum temperature on the TiC surface and the linear interpolation value from the

100 11 80

10 o

/

>

/ /

,,..z 6 0 /;

< z

,

/

9

8

P=1.2 ,'

4O

/

# //

6

D

t-

5

O 20 0

0

4

/

et al.

measured temperature on the Cu side were adopted. Other properties of the FGM layers were estimated based on the experimental values and their linear interpolation.

3 RESULTS AND DISCUSSION The temperature distribution on the surface of the TiC side of the WU-80 specimen measured in the steady state is shown in Figure 4, which indicates that the temperature changed gradually from the center (1250K) to the edge (850 K). As shown in Table 2, the maximum of the TiC side surface temperature (Ts) of the WU-80 specimen increased by 50 K after the 20th thermal shock test and reached 1250 K; while the WU-81 and 82 specimens showed temperature increases of 200 and 150 K, respectively. The temperature change on the TiC surface depends on the microstructural variation of the specimens due to the cyclic thermal shock tests because the bottom surfaces of all specimens welded to the Cu plate were steadily cooled with liquid nitrogen. In thermal shock tests, the WU-81 specimen showed a large change in structure; in contrast, the WU-80 specimen revealed only a small change. The surface temperature, Ts (TiC side), bottom temperature, Tb (Ni3A1 side) and their difference during the thermal shock tests for the WU-80 specimen are shown in Figure 5. There was only a small change in the maximum surface temperature of TiC during each

I

0

1

2

3

4

5

6

Distance from surface of TIC side, x / mm Figure 2 Relationship between the compositional distribution and

thickness for WU-80

Radiationpyrometer

Video camera

Vacuum vessel

Reflector

__n,

Liq. N2 tank

Test-board Sample uum pump

~

Quarlzglass

,~ ~ DC power supply - - I

I~

ill I 1 ~ I I

~'Cut-off valve "Shutter ~--Refleclor

Xearclamp

_1 Figure 3 Schematic diagram of test apparatus creating great tem-

perature difference within a sample 3

Figure 4 Temperature distribution on the surface of TiC layer of

WU-80 Table 2 Change of maximum temperature on the TiC surface of the

specimens with cyclic thermal shock tests Temperature (K) Number of cyclic thermal shock tests

WU-80

WU-81

WU-82

1 20

1200 1250

1030 1230

1310 1460

23

Cyclic thermal shock of TiC/Ni3AI FGMs: L. M. Zhang thermal shock test. The temperatures at the Ni3A1 side also increased correspondingly, the temperature differences (Ts-Tb) in the specimen showed little change, remaining at a relatively stable level (between the a-b line and the c-d line in Figure 5). Moreover, the temperature change as shown by the a - d line increased only about 10% during the cyclic thermal shock process. The effective thermal conductivities of the specimens, Kerf (Keff = qt/(Ts-Tb), where q is the heat flux and t is the thickness of a specimen) were calculated from the measured Ts-Tb values. The effective thermal conductivities, Kerf, and the conductivities normalized by the value of the first cycle, nk, are shown in Figure 6 as a function of the number of cyclic thermal shock tests. After the 20th thermal shock, the effective thermal conductivities of WU-80, 81, and 82 decreased to 82.2%, 47.6%, and 59.7% of their initial values, respectively. The changes in Kefr of WU-80 displayed a slow degradation mode l°. The changes in K~fr of WU-81 and 82 displayed two degradation stages, i.e. Keff values of WU-81 and WU82 steeply decreased during the early thermal shock cycles, then slowly degraded like that of WU-80. It is known that the Kerr variation of WU-80 is in agreement with slow degradation mode of Kerr of FGM specimens, i.e. the damage takes place on a microstructure scale. WU-81 and 82 conform to a medium degradation mode, i.e. microstructural damage such as microcracks on the surface or in the interior were observed in the primary stage, such damage resulting in a relaxation of thermal stress by the energy dissipation effect due to the friction of microcrack boundariesil. After the cyclic thermal shock tests, small cracks were detected near the TiC side in WU-80 as shown in Figure 7(a). Visible cracks were observed in WU-81 (Figure 7(b)) and WU-82, with the cracks on the TiC surface of WU-81 being comparatively larger. However, none of the specimens were damaged in the thickness direction, because the cracks caused by the Orr and ~r00stresses were vertical to the TiC surface. The vertical cross-sections of specimens after the thermal shock tests are shown in Figure 8. The cracks on the surface of the TiC layer can be seen to initiate in the vertical direction, then propagate toward the interior. As shown in Figures 8(b) and (c), vertical cracks in WU-81 and 82 evolve into horizontal cracks, while there are no horizontal cracks in WU-80 (Figure 8(a)) even though the vertical cracks in this specimen change directions twice. The above observations, combined with the results of Figure 6, suggest that the steep decrease in Kerf was due to the horizontal extension of the vertical cracks. Vickers hardness distribution in the thickness direction was measured on the cross-sectional surface. Hardness values at positions, C, R, and L of WU-80 are shown in Figure 9. It was found that the hardness decreased with the increase of Ni3A1 content. At the position C, there was an obvious hardness decrease on the TiC surface. The decrease in the hardness at position C was larger than the corresponding decrease at positions L and R. Moreover, the hardness decrease at the

24

et al.

D*tt*RUN(WU-80)

14110 . . . . . . . . .

1200

2000

T:II

m T I - T

b (K)

1500

iooo 800 - -

1000

! •...,/"

2OO 0

,

,

s~

~ ..z.../,,: .;\,;~"J:'/"\:"-/::'%",i:;',,:',./

,

,

0

I

,

i

i

,

I

100

. . . .

I

200

. . . .

I

300

. . . .

400

I

. . . .

500

E E

0

600

Time of the thermal shock tests, t / sec

Figure 5 Surface temperature (Ts: TiC side), bottom temperature (Tb: NiaAI side) and their differences (Ts-Tb) with time of a thermal shock test for WU-80

30

~

o .e.

n~(WU-eo) | r~(WU-O~) -t 1.0

0, i;

20

,==, lO I

0

i

~= o

K~ (WU,.80) K~ (WU-81 ) K~ (WU'82)

~. 41' i

i

I

I

5

Z i

i

i

t

I

10

i

i

i

i

I

i

I

15

k

i

I

0

20

Number of cyclic thermal shock testa

Figure 6 Effect of cycles of thermal shock tests on the effective thermal conductivities of TiC/Ni3A1 FGMs

Figure 7 Surface of TiC sides of the TiC/Ni3AI FGMs after the thermal shock tests: (a) WU-80 and (b) WU-81. Arrows indicate cracks

Cycfic thermal shock of TiC/Ni3AI FGMs: L. M. Zhang et al.

. 25 I

-

15

~

10

.m ~ u

s

m o Q

TiC

C

~ L .....

~1 20

R

Ni~J Section of FGM

C W

0

0

1000

t

2000

3000

4000

5000

6000~ 7000 t

Distance from TIC side, x / p.m

TIC

N 13AI

Figure 9 Vickers hardness on the cross-sectional surface of WU-80 after the cyclic thermal shock tests

20 m a.

i

,

' Ni~I

~15 > .

|10 C

,IC

5 _o

i

0o TIC

,

,

,

i

. . . .

1000

i

. . . .

2000

i

. . . .

3000

i

. . . .

4000

i

. . . .

5000

Distance from TIC side, x / p m

I,,-,

,

6000

t

7000

NO=AI

Figure 10 Vickers hardness at the position C on the cross-sectional surface for TiC/Ni3AI FGMs after the cyclic thermal shock tests

Cross-sectional view of the TiC/Ni3AI FGMs after the cyclic thermal shock tests. (a) WU-80 (the arrows show change in the direction of the cracks), (b) WU-81 (the arrows show a horizontal extension of the crack) and (c) WU-82 (the arrows show the extension of the cracks) Figure 8

position C of WU-82 was larger than that of WU-80 and 81 as shown in Figure 10. The lower hardness near the TiC layer for WU-82 was due to the presence of porous structures. Figure 11 shows the mesh employed in the finite element analysis (A), and contour plots of the three main stresses: through-thickness stress a= (B); radial stress O'rr (C); and hoop stress ~00 (D), for the steady state of WU80. The maximum through-thickness tensile stress was about 2 GPa at the center of the restrained edge of the Cu plate. The Cu plate did not separate from the sampleholder due to its shape deformation. Further, another peak of through-thickness tensile stress in the layers of

the specimen was about 500 MPa, not large enough to separate the specimen because the layers had greater strength. There was also a tensile stress peak of 100 MPa at the TiC side and the free edge, which was not large enough to cause delamination. The radial stress, arr, and the hoop stress, a00, appear to be very important. On the TiC surface and the parallel layers 2 mm beneath the TiC surface, the maximum radial stress, O'rrmax, and the maximum hoop stress, a00max, were - 1 . 6 and - 1 . 5 GPa, respectively, around the zz-axis. Both O'rr and or00stresses decreased in the radial direction. On the TiC surface, a state of zero stress appeared at a position 2/3 of the radius in the radial direction. From the analysis mentioned above and the results of microstructural observation and hardness measurement, it is suggested that the large compressive inplane stresses in the pure TiC layer and the layers nearby led to the formation of cracks. However, according to the elastic theory, if the TiC does not generate plastic deformation, the stress will return to the initial state after cooling and crack formation will be avoided. Studies on plastic properties of TiC at high temperatures indicate that ductileto-brittle transition temperatures of TiC are between 1073 and 1173 K and that its compressive yield strength is about 1 GPa ]2']3. From the calculated results of thermal stresses based on the experimental values in the steady state, it is known that compressive inplane stresses

25

Figure 11

(A) : Mesh geometry employed In the ¢xisymmetrlof-e analysis

(B) : ~u (P" 1.2)

(C) : err (P = 1.2)

(D) : e00 (P " 1.2)

Mesh geometry and distribution of thermal stresses under the steady-heating condition for WU-80

strength of TiC materials. Figure 12 indicates that P = 1.2 for the thermal stress simulation of WU-80 corresponds to the optimum design for the residual thermal stress 5. Within the exponent range considered, the radial and hoop stresses decrease to the minimum values, which indicates that the change of distribution exponent cannot decrease the compressive stress of WU-80 any more. The only method ,gf improving the surface stress state is thought to be the adjustment of the thickness-diameter ratio of disk-shaped FGM specimens. This will be the focus of future studies.

1000 m L

500 o

o

E

--D..------tS--• • ,L

-500

W

-lOOO

On..ten" (Y0O-ten (~zz-ten. (Yrr-com (~00-com. (Yzz-eom.

E = -1500 -2000 -2500

i

0

0.4

L

L

I

I

I

0.8 1.2 1.6 2 Distribution exponent, P

I

I

2.4

2.8

Figure 12 Relationship between the m a x i m u m stress and distribution exponent under the steady-heating condition for WU-80

appeared in the region 0 ~ 2/3 of the radius in the radial direction, and that their magnitudes were in the range of -1.6GPa to -500MPa when the TiC surface was exposed to a temperature of 1208K. The maximum compressive in-plane stresses in the center can possibly result not only in plastic deformation but also in compressive damage before the cooling phase commences. Therefore, the tensile stress generated on the surface during the cooling phase can produce vertical cracks and can also propagate cracks which already exist as shown in Figure 8. It is concluded that the driving force for the formation of vertical cracks was the excessive compressive stress in the TiC layer. It is very important to reduce the driving force so that the compressive stress is less than the compressive yield

26

4 CONCLUSIONS (1) Among TiC/Ni3A1 FGM specimens with the same densified structure, a larger ratio of thickness to diameter leads to better cyclic thermal shock resistance. Porous TiC and TiC-rich layers are beneficial for thermal stress relaxation, but crack formation cannot be avoided due to their low strength. (2) There is a close relationship between the change of effective thermal conductivity and the features of crack-propagation in FGMs. The steep degradation of effective thermal conductivity is mainly caused by the horizontal extension of vertical cracks. (3) The curve of the normalized effective thermal conductivity variation shows that there are two modes of damage. One is the slow degradation mode which appeared in the WU-80 specimen, and the other is the crack extension and energy dissipation mode which appeared in the WU-81 and 82 specimens. (4) Crack formation in FGMs is mainly caused by the excessive compressive stress on the TiC surface in a

Cyclic thermal shock of TiC/Ni3AI FGMs: L. M. Zhang steady state. Controlling the compressive stress below the yield strength of TiC is the key for preventing the vertical crack formation.

4 5 6

ACKNOWLEDGEMENTS

7 8

This work was supported by the Japan 'JSPS' and Chinese '863' Programs. 9

REFERENCES 1 2 3

Niino, M., Hirai, T. and Watanabe, R. J. Japan Sot:. of Comp. Mater. 1987, 13, 257-261 Sasaki, M., Hirai, T., Hashida, T. and Takahashi, H. J. Japan Soc. of Powder and Powder Metallurgy 1990, 37, 966 967 Kumakawa, A., Niino, M., Kiyoto, S. and Nagata, S. "Ceramic Transactions Vol. 34: Functionally Gradient Materials', (Eds J.B. Holt, M. Koizumi, T. Hirai, Z.A. Munir), Am. Ceram. Soc., Westerville, Ohio, USA, 1993, pp. 213-220

l0

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et al.

Kawasaki, A., Hibino, A. and Watanabe, R. J. Japan Inst. Metals 1992, 56, 472-480 Zhang, L.M., Liu, J., Yuan, R.Z. and Hirai, T. Mater. Sci. & Eng. A 1995, A203]1-2, 272 277 Mattox, D.M., Mullendore, A.W., Pierson, H.O. and Sharp, D.J.J. Nucl. Mater. 1979, 85-86, 1127-1131 Motoki, Y. 'Ceramics Engineering Handbook', Japan Ceram. Soc., Gihoudou Press, 1989, pp. 2102-2113 Zhang, L.M., Omori, M., Yuan, R.Z. and Hirai, T. 'Proc. of 3rd. Int. Sym. on Structural and FGM', (Eds B. Ilschner and N. Cherradi), Presses Polytechniques et Universitaires Romandes, Lausanne, Switzerland, 1995, pp. 59 64 Zhang, L.M., Yuan, R.Z. and Hirai, T. J. Mater. Sci. Lett. 1995, 14, 1620 1623 Kumakawa, A., Takahashi, M., Sasaki, M., Togawa, M., Kitaguchi, S. and Nishimori, H. 'Proc. of 3rd. Int. Sym. on Structural and FGM', (Eds B. llsehner and N. Cherradi), Presses Polytechniques et Universitaires Romandes, Lausanne, Switzerland, 1995, pp. 391-396 Lee, W.J. and Case, E.D. Mater. Sci. & Eng. 1989, All9, 113 126 Das, G., Mazdiyasni, K.S. and Lipsitt, H.A.J. Am. Ceram. Soc. 1982,65, 104 110 Miracle, D.B. and Lipsitt, H.A.J. Am. Ceram. Sot'. 1983, 66, 592 597

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