Creep behavior of a reinforced Al-7005 alloy: Implications for the creep processes in metal matrix composites

Acta mater. Vol. 46, No. 4, pp. 1143-l 155, 1998 0 1998 Acta MetallurgicaInc. Published bv Elsevier ScienceLtd. All rights reserved Printed in &eat Britain PII: S1359-6454(97)00320-O 1359-6454198 $19.00+ 0.00

CREEP BEHAVIOR OF A REINFORCED Al-7005 ALLOY: IMPLICATIONS FOR THE CREEP PROCESSES IN METAL MATRIX COMPOSITES YONG

LI and T. G. LANGDON

Departments of Materials Science and Mechanical Engineering, University of Southern California, Los Angeles, CA 90089-1453, U.S.A. (Received

12 May 1997; accepted 26 August 1997)

Abstract-Creep tests were conducted on an Al-7005 matrix alloy reinforced with 20 vol.% of irregularly shaped A1203 particulates. The results, which cover almost six orders of magnitude of strain rate, give high values for the apparent stress exponent and the apparent activation energy for creep. The introduction of a threshold stress into the analysis leads to a stress exponent of -4.4 and data which are consistent with creep controlled by a dislocation climb process and class M behavior in the Al-7005 matrix alloy. These results are compared with an earlier report of creep behavior in an Al-6061 matrix composite, also reinforced with 20 vol.% of irregularly shaped Al203 particulates, where the stress exponent was close to three and the data were consistent with control by a dislocation glide process and class A behavior in the Al-6061 matrix alloy. It is shown that the creep of metal matrix composites is controlled by creep of the matrix alloys and there is a consequent division of the creep behavior of the composites into two types, class M and class A, as in the creep of solid solution alloys. cl 1998 Acta Metallurgica Inc.

1. INTRODUCTION

An earlier report described the creep behavior of an Al-6061 metal matrix composite reinforced with 20 vol. % of irregularly shaped A&O3 particulates and fabricated using an ingot metallurgy procedure [hereafter designated Al 6061-20 vol.%AlzOs(p) where p denotes particulates] [l]. The creep data, in the form of a logarithmic plot of strain rate vs stress, extended over five orders of magnitude of strain rate and exhibited the characteristic curvature indicative of the presence of a threshold stress for creep. By incorporating a threshold stress into the analysis, it was shown that the true stress exponent for creep, n, was close to three and the true activation energy for creep, Q, was similar to the value anticipated for diffusion of Mg in the Al-based lattice. This result suggests, therefore, that the ratecontrolling mechanism during creep of this composite is, as in many dilute Al-Mg solid solution alloys [2,3], the viscous glide of dislocations associated with the dragging of Mg atom atmospheres. In addition, this interpretation is consistent also with creep data obtained on an Al-6061 alloy reinforced with essentially spherical alumina-based microspheres [4] and with a recent re-interpretation of published creep data for an Al-4% Mg composite reinforced with 10 vol.% of Sic particles [S]. These creep results are significant in two respects. First, they serve to demonstrate that the creep of metal matrix composites is controlled by the rate of deformation within the matrix alloy. Second, they provide the first direct evidence for the possibility 1143

of control by viscous glide in a metal matrix composite. It is well documented that the creep behavior of solid solution alloys is divisible into two distinct types, designated class M (metal type) and class A (alloy type) [6]. In class M, which is similar to pure metals, dislocation climb is the rate-controlling process with a stress exponent close to five and an activation energy similar to the value for lattice selfdiffusion; whereas in class A, which applies in some solid solution alloys under well-defined conditions, viscous glide is rate-controlling with a stress exponent equal to three and an activation energy associated with interdiffusion of the solute atoms. There is also generally a transition from class M behavior at lower stresses to class A behavior at higher stresses with the level of the transition stress dependent upon the testing temperature, the concentration of the solute, the size difference between the solute and the solvent atoms and the stacking fault energy of the material [3]. If indeed the creep of metal matrix composites is controlled by deformation of the matrix alloy, there is the clear implication that the creep behavior of these materials may divide into two separate classes equivalent to the divisions into class A and class M already documented in some detail in solid solution alloys. The present investigation was motivated by this possibility. The standard AI-6061 matrix alloy contains -1 wt% Mg as the major alloying element and it is well established that dilute AI-Mg alloys exhibit

LI and LANGDON:

1144

Table 1. Nominal Alloy Al-6061 Al-7005

Si

Fe

0.4-0.8 0.35

0.7 0.4

CREEP BEHAVIOR OF 7005 ALUMINUM ALLOY

compositions CU

O.lS-0.40 0.10

in the Al-6061 and AI-7005 matrix alloys (wt%) Mn

Mg

Cr

Zn

Ti

Al

0.15 0.2-0.7

0.8-1.2 1.0-1.8

0.04-0.35 0.06-0.20

0.25 4.0-5.0

0.15 0.01-0.06

Bal. Bal.

extensive creep behavior within the class A region [6-S]. By contrast, the Al-7005 alloy contains -4.5 wt% Zn as the major alloying element and experiments have established that dilute Al-Zn alloys tend to exhibit class M behavior over a very wide range of stresses with a stress exponent of n = 4.4-4.6 [9-121. Therefore, the objective of this investigation was to determine the creep characteristics of a metal matrix composite based on Al-7005 as the matrix alloy. As will be demonstrated, the results provide support for the concept of two separate classes of creep behavior in metal matrix composites.

2. EXPERIMENTAL MATERIALS AND PROCEDURES The investigation was conducted using an Al7005 alloy reinforced with 20 vol.% of irregularly shaped A1203 particulates [henceforth designated Al 7005-20vol.%AlzO&)]. The composite was fabricated using a proprietary casting technique and it was supplied by Duralcan U.S.A. (San Diego, California) in the form of rods with a diameter of 19.1 mm. As with the Al 6061-20vol.%A120~(p) tested earlier [l], the particles ranged in size from -16 to -23 pm with an average size of -20 pm. Table 1 gives the nominal composition of the Al7005 matrix alloy and also provides, for reference, information on the Al-6061 alloy. Tests were conducted under double-shear conditions [13] using the same specimen configuration as documented in other reports [l, 141. A modified T53 heat treatment is standard for the Al7005 alloy and consists of a solution treatment at 753 K for 2 h, water quenching, followed by two ageing treatment at 378 K for 8 h and subsequently at 423 K for 12 h. All specimens were subjected to this modified T53 heat treatment prior to testing. Figure 1 shows the morphology and distribution of A1203 particulates in the Al-7005 matrix after the heat treatment and it is apparent from inspection that the alumina particulates are distributed essentially uniformly throughout the matrix. The mean linear intercept grain size of this material was estimated as -20 pm. The specimens were creep tested in air under conditions of constant shear stress over a temperature range from 573 to 773 K and with the testing temperature continuously monitored and maintained constant to within f2 K of the desired value. Shear strains were measured using a linear variable differential transducer and they were recorded continu-

ously on a strip chart recorder. Tests were conducted over a range of stresses to give measured strain rates covering almost six orders of magnitude. Following creep testing, samples were prepared for examination by transmission electron microscopy (TEM). Slices with a thickness of -500 pm were cut from within the gauge lengths of selected specimens and these slices were polished mechanically to a thickness of -30 pm and then thinned to perforation in a twin-jet electropolishing unit using a mixture of 25% HNOs and 75% methanol maintained at a temperature below 240 K with dry ice. Observations were performed using a CM-20 Philips transmission electron microscope with an operating voltage of 200 kV.

3. EXPERIMENTAL RESULTS 3.1. Creep curves Typical creep curves are shown in Fig. 2 for a testing temperature, T, of 573 K, plotted in the form of the shear strain, y, vs time, t, for different values of the shear stress, r. These creep curves, and others obtained under different testing conditions, show the typical characteristics of creep behavior in metal matrix composites: namely, a short primary stage, a very brief secondary stage representing a minimum in the measured creep rate and then an extended tertiary stage [l, 4,151. The three creep curves shown in Fig. 2 are replotted in Fig. 3 in the form of the shear strain rate, Jo,vs the shear strain, y, thereby demonstrating the lack of an extended period of steady-state creep.

Fig.

1. Microstructure of the Al 7005-20vol.%A1203(p) composite after the modified T53 heat treatment.

LI and

LANGDON:

CREEP BEHAVIOR OF 7005 ALUMINUM ALLOY

IO-’ m I Al 7005-2Ovol.% $0,(p)

Al 7005-2Ov0l%~0,@)

10’

rw

~

/en

1

1145

0

12.0

0

16.2

A

24.7

-

o-_o-o-o-o1

2

I

h

7

.e

t @)

IO-5

Fig. 2. Typical creep curves for a temperature of 573 K 3.2. Stress and temperature dependence of the minimum creep rate All of the creep data obtained in this investigation, for four testing temperatures from 573 to 773 K, are plotted logarithmically in Fig. 4 in the form of the minimum creep rate vs the shear stress. These creep data span almost six orders of magnitude of strain rate and they show the characteristic curvature indicative of the presence of a threshold stress and now generally associated with the creep of metal matrix composites when data are collected over a very wide range of strain rates [l, 15-231. Because of the curvature in Fig. 4, the values of the apparent stress exponent, n,, and the apparent activation energy, Qa, are functions of the applied stress, and calculations show these values lie within the ranges of n,=6-17 and Q,=170-380 kJmol_‘, respectively. These values are very high by comparison with the stress exponents of n s 4.4-4.6 reported for the creep of Al-Zn solid solution alloys [9-121 and with the anticipated activation energy for creep which should be close to the value for self-diffusion in Al (-143.4 kJ mol-’ [24]). 3.3. Microstructural observations after creep testing Careful inspection by TEM revealed the presence of many precipitates having a reasonably uniform distribution: an example is shown in Fig. 5(a) for a

1 10’

102

T Wa) Fig. 4. Minimum creep rate vs shear stress for tests conducted at temperatures from 573 to 773 K. specimen tested at a temperature of 623 K under a shear stress of 20.5 MPa. These precipitates were generally small and essentially spherical, with sizes in the range of -15-30 nm, but there were also some larger precipitates having sizes up to a maximum of -100 nm. It was found by energy dispersion spectroscopy that some of these precipitates contained an excess of Cu, suggesting the presence of semi-coherent MgCu2 particles. There was also direct evidence for interactions between the precipitates and dislocations: these interactions are apparent in Fig. 5(a) and a clear example of a dislocation attached to precipitates is shown in Fig. 5(b).

4.

DISCUSSION

4.1. Creep properties of Al 7005-20vol.%Al,O,(pj IO”

0

104

162 A ? ? 247

-j

5

F-

120

T=573K

\ A ?-AA

AA’

\

0’

lo-’

Fig. 3. Shear strain rate vs shear strain for the three specimens shown in Fig. 2.

The values of the individual threshold stresses may be estimated at each temperature by using a linear extrapolation procedure in which i)‘!” is plotted against z on linear axes for selected values of n and, in the absence of any significant curvature, the plots are extrapolated linearly to intersect the stress axis at zero strain rate [25]. This procedure was followed in the present investigation using values of n of 3, 4.4 and 8. These values were selected because the viscous glide mechanism requires n = 3 [2,3] and this value was found to be appropriate for the Al 6061-20vol.%Alz0,(p) composite [l], a value of n = 4.4 is appropriate for

1146

LI and LANGDON:

CREEP BEHAVIOR OF 7005 ALUMINUM ALLOY stress decreases with increasing testing temperature, as reported also in numerous investigations of metal matrix composites [l, 15,19,29,30], unreinforced alloys [3 l-351 and dispersion strengthened alloys [3641]. Taking the values for r. listed in Table 2, Fig. 7 shows a logarithmic plot of the minimum creep rate vs the effective stress, (r - ro). Inspection shows that, except only for the experiment conducted at the lowest level of stress at the lowest testing temperature, all of the datum points fit well to a series of parallel lines having a slope which gives n = 4.4. From this analysis, therefore, the creep rate for the Al 7005-20vol.%A1~03(p) composite can be expressed in the form ADGb 9=F

Fig. 5. Transmission electron micrographs showing (a) precipitates in a specimen tested at a temperature of 623 K under a shear stress of 20.5 MPA and (b) an example of dislocation attachment to precipitates in the same specimen.

creep data for Al-Zn solid solution alloys [9-121 and, in addition, it is the value reported for pure Al from an extensive compilation of creep data [26] and a value of n = 8 is associated with a constant structure model of creep [27,28]. These three plots are shown in Fig. 6(a)-(c) and inspection shows there is significant curvature in the plots constructed with values of n of 3 and 8 whereas the plot with n = 4.4 shows a good linear correlation. This result is particularly striking when it is noted that the earlier analysis of creep data for the Al 6061-20vol.%Al~O~(p) composite revealed linearity with n = 3 and marked curvature when using values of 12 of 5 or 8 [l]. Using the data plotted in Fig. 6(b) with R = 4.4, the straight lines may be extrapolated to zero strain rate to give the values for the threshold stress, zo, which are listed in Table 2. These results show that the threshold

Z-T~ n (7 >

(1)

where D is the appropriate diffusion coefficient [ = D,, exp (--Q/K’), where Do is a frequency factor and R is the gas constant], G is the shear modulus, b is the Burgers vector, k is Boltzmann’s constant, A is a dimensionless constant and the value of n is equal to 4.4. It follows from equation (1) that the value of the true activation energy for creep may be estimated from a semi-logarithmic plot of 9G” - ‘T against l/T. Taking n = 4.4 from Fig. 7 and with (z - ro) = 3 MPa, this plot is given in Fig. 8 and it leads to a value of Q ‘n 120 kJ mol-‘. This value of Q is slightly lower than the activation energy for self-diffusion in pure Al (143.4 kJ mol-’ [23]) but it is similar to the activation energy of 128 f 3 kJ mol-’ obtained experimentally for creep in a region where n = 4.5 in an Al-Zn solid solution alloy containing 10 at.% of Zn (equivalent to 21.2 wt% Zn) [12]. In an investigation of creep in an Al-10 at.% Zn alloy by Soliman and Mohamed [12], an extensive and Q= 128 i 3 creep regime having n-4.5 kJ mol-’ was interpreted, based on the shapes of the creep curves and the nature of the transients in stress reduction experiments, in terms of class M creep behavior and control by a dislocation climb process. The similarity in the values of n and Q obtained in this investigation suggests that it may be appropriate to adopt the same approach in analyzing the creep data for the Al 700520vol.%Al~O~(p) composite. If creep is controlled by dislocation climb, the diffusion coefficient for the climb process, D,, replaces D in equation (1). Therefore, all of the data may be normalized through a logarithmic plot of jkT/DcGb vs (z - 70)/G. In order to calculate the appropriate value for D,, it is necessary to use the approach developed by Fuentes-Samaniego et al. [41,42] to estimate the coefficients, 0,” and Di, associated with diffusion-controlled climb and glide creep processes in binary solid solution alloys. This approach is examined in Appendix A with reference

LI and LANGDON:

CREEP

BEHAVIOR

OF 7005 ALUMINUM

ALLOY

1147

(a> 0.20

1

I

Al 7005-20~01.

I

I

I

0.4 Jb)

I

I

I

% AJO,

I

I

Al 7005-20~01.

0.15

I

%Al,O,

,

I

I

(p)

0.3

h d: ; fi ; :?-

27 cf5 0.10 P ?i=-

0.2

0.05

rw

0.1 0

573

0

623

A

673

v

0.00

0.0 0

10

20

30

0

40

10

773

L

I

20

30

r WW

T WW C

0.6

( )

,

I

Al 7005-2Ovol.

I

I

% Al,O,

I

I

I

(p)

t 0.5

0.4

Tu

0

10

20

0

573

0

623

A

673

v

773

30

40

t @@a) Fig. 6. The linear extrapolation

method

for estimating the threshold (b) 4.4 and (c) 8.

i

stresses using values for n of (a) 3,

I

40

LI and LANGDON:

1148 Table

2.

Values

of

CREEP BEHAVIOR OF 7005 ALUMINUM ALLOY

the threshold stresses 20vol.%Al~O~(p)

in

Al

7005-

TO (MW

T (K) 573 623 673 773

7.8 5.1 3.6 1.3

to the Al-Zn system and with the solute concentration, c, taken as 0.021 (equivalent to 4.5 wt% Zn). The individual values of Da are estimated in Appendix A for each of the testing temperatures used in this investigation and these values are listed in Table Al. Putting D, equal to D,“, Fig. 9 shows a normalized plot of the data in which all of the points, except only for the single point at the lowest stress level at 573 K, are brought together onto a single line with a slope of n = 4.4. 4.2. Significance

I

v

I

V

I

It is apparent from Table 2 that the estimated threshold stresses decrease with increasing temperature. The normalized threshold stress in the creep of metal matrix composites is often expressed by a relationship of the form [ 151

II/

s=Baexp

0 0

A

I/ /

I

;

4.4

?? caI0

?I 1

vB’ / 0 Ic7 1.

-

Fig. 7. Minimum creep rate vs the effective shear stress for tests conducted at temperatures from 573 to 773 K.

10’0 L 1.2

I

I 1.3

& (

A I

of the threshold stress

I

I 1.4

I

(2)

>

where Qc is an energy term associated with the process which gives rise to the threshold stress and B. is a constant. By constructing a semi-logarithmic plot of rc/G against the reciprocal of the absolute temperature, as shown in Fig, 10, the value of Qc for the Al 7005-20vol.%A1~03(p) composite may be estimated as -27 kJ mol-‘. The precise significance of Qc is not fully understood at the present time. Nevertheless it is important to note that several experimental values of Q. are now available, covering a range of Al-based materials, and all of these values lie within the relatively narrow range of -19-28 kJ mol-i. This may be seen by inspection of the data given in Table 3 which cover four different metal matrix composites fabricated by powder metallurgy [15,16,21] and

I

I

1.5

I

1.6

I

I

1

1.7

1.8

1000/T (K-l) Fig, 8. Semi-logarithmic

plot of ~JG~.~Tvs l/T in order to obtain for creep, Q.

the value of the true activation

energy

LI and LANGDON:

CREEP BEHAVIOR OF 7005 ALUMINUM

ALLOY

1149

Table 3. Experimental values reported for Q0 in equation (2)

10-6

Material

Al 7005-2Ovol. % GO3 (p)

0’

Tw 0 0

573 623

Q,, (kJ mol-‘)

Al 700520vol.%Al,01(p) Al 6061-3Ovol.%SiC(p) Al-30vol.%SiC(p) Al 6061-20vol.%AlZ03(p) PM 6061 Al alloy? PM 2024 Al alloy PM 2124 Al alloy RS Al-Si&Ni-Cr alloyt

J

-27 -19.3 -23 -25 -19.3

Reference This investigation U&161

1211 [II

-23

[311 [351

-27.5 -21

[491 [341

fPM = powder metallurgy; $RS = rapid solidification

Fig. 9. Normalized shear strain rate vs normalized effective shear stress for tests conducted at temperatures from 573 to 773 K. liquid metallurgy [l] procedures and including both A1203 and SIC reinforcements, three Al alloys fabricated by powder metallurgy [31,35,49] and an Al alloy fabricated by rapid solidification [34]. Therefore, it is concluded that the value of Qa is relatively insensitive both to the nature and to the

volume fraction of the reinforcement in metal matrix composites and similar values of Qa are obtained in Al-based alloys without any reinforcement. It is instructive to note also that the range of experimental values for Qo, which corresponds to -0.2-0.3 eV, is similar to the anticipated values for the binding energies between dislocations and impurity atoms in the glide plane [50]. In detailed investigations of the creep behavior of an Al-6061 matrix alloy containing 30 vol.% of Sic particles [15] and an unreinforced Al-6061 alloy [31], both prepared using powder metallurgy procedures, the threshold stresses were attributed to an interaction between mobile dislocations and an array of fine incoherent oxide particles introduced during atomization in the fabrication process. Alternatively, the threshold stresses observed in the creep of an Al 6061-20vol.%Al~O~(p) composite fabricated by a liquid metallurgy process were attributed both to the presence of many fine spine1 particles which are known to be introduced by oxidation of the melt during processing and to the transformation after ageing of p precipitates into incoherent platelets of equilibrium P-Mg$i precipitates [ 11.

AI 7005-2oVoI%AI,O,(p)

IO-5 1.2

1.6

1.4

1.8

1000/T (K-l) Fig. 10. Semi-logarithmic plot of the normalized threshold stress vs the reciprocal of the absolute temperature.

1150

LI and LANGDON:

CREEP BEHAVIOR

In order to determine the origin of the threshold stresses in the Al 700520vol.%Al~Os(p) composite, it is important to note that Al-Zn and Al-Zn-Mg alloys are known to exhibit complex sequences of precipitation [51-531 and Al-Zn-Mg alloys are strengthened by precipitation reactions accompanying the ageing treatments [54]. For example, in Albased alloys having Mg/Zn ratios of 1:2 to 1:3, as in the Al-7005 matrix alloy used in the present investigation, it has been established that binary MgZnz precipitates are nucleated which are stable at relatively low temperatures [51], and recent experiments on a commercial Al-Zn-Mg alloy revealed the nucleation of semi-coherent Mg(Zn,Al,Cu)z precipitates having nanometer sizes on ageing at a temperature of 423 K [55]. These precipitates represent effective obstacles to dislocation motion and they provide, together with any MgzSi precipitates which may be present in the matrix alloy [56], a probable source for the threshold stresses recorded in the present investigation. Furthermore, the TEM observations recorded in Section 3.3 confirm both the presence of extensive precipitation [Fig. 5(a)] and the occurrence of dislocation-precipitate interactions [Fig. 5(b)]. The possible origin of the mechanism giving rise to the threshold stresses is examined in Appendix B. 4.3. A comparison

of class M and class A creep behavior in metal matrix composites

The results obtained in the present investigation on the Al 7005-20vol.%Al~Os(p) composite are remarkably different from those reported earlier for an Al 6061-20vol.%Al~0,(p) composite [l] despite the fact that both sets of experiments were conducted under double-shear conditions and over similar ranges of temperatures. In both investigations, curvatures in the basic logarithmic plots of shear strain rate against shear stress were interpreted by introducing a threshold stress into the analysis, but the Al 7005-20vol.%Al~O~(p) composite gives clear evidence for class M behavior with n = 4.4 and Q = 120 kJ mol-’ and the Al 606120vol.%A120s(p) composite provided strong support for control by class A behavior with II r: 3 and Q ‘c 130 kJ mol-’ [l]. These well-defined differences are especially striking when it is noted that (i) both composites were obtained from the same supplier and they were both fabricated using the same ingot metallurgy procedure, and (ii) both composites contain the same reinforcement in terms of type (A120s), size (-20 pm) and volume fraction (20%). Therefore, the marked variation in the creep behavior must be due unambiguously to the difference in the matrix alloys. It is apparent from Table 1 that the major alloying elements are -4.5 at% Zn in the Al-7005 alloy

OF 7005 ALUMINUM

ALLOY

and -1.0 wt% Mg in the Al-6061 alloy. By comparison with creep investigations of the unreinforced Al-Zn [9-121 and Al-Mg [6-81 solid solution alloys, it is reasonable to conclude that class M behavior and control by dislocation climb will be of major importance over a wide stress range in the Al-7005 alloy whereas class A behavior and control by dislocation glide will be of major importance in the Al-6061 alloy. Although these anticipated trends are consistent, at least quantitatively, with the experimental data for the two metal matrix composites, a clear distinction between class M and class A creep behavior necessitates a more detailed analysis. First, inspection of Table 1 shows that the Al7005 matrix alloy contains -1.4 wt% Mg in addition to -4.5 wt% Zn and yet the presence of Mg atoms in the Al-7005 alloy has not produced class A behavior in this composite. This effect may be understood by noting that the formation of precipitates, such as MgZn2 and Mg$i, leads to a depletion of Mg within the matrix alloy, as documented in an examination of an Al-Zn-MgCu alloy composite with SIC reinforcement [56]. Second, any interpretation of the behavior of metal matrix composites in terms of the two classes of creep already established for solid solution alloys requires a demonstration that the creep data for the composites are consistent with the prediction which has been developed to govern the advent of class M and class A behavior in the unreinforced alloys. This analysis can be undertaken by using the criterion developed earlier to predict the transition from class M behavior at the lower stresses to class A behavior at the higher stresses [3]. Specifically, it was shown earlier that the transition occurs at a well-defined experimental condition which depends upon several factors including the concentration of the solute and the size of the solute atoms with respect to the atoms of the solvent. However, the transition developed for solid solution alloys must be modified to incorporate both the possibility of different viscous drag mechanisms occurring within the class A region and, for composite materials, the replacement of the applied stress with an effective stress as in equation (1). The method of developing and utilizing a relationship marking the transition between class M and class A creep behavior in metal matrix composites is outlined in Appendix C and the transition is represented by equation (ClO). This transition may be depicted graphically and the result is shown in Fig. 11 where control by climb and class M is dominant on the left, control by glide and class A is dominant on the right, and the boundary between these two regions was established earlier for an Al3% Mg alloy [3] but was shown subsequently to be equally applicable to Al-Zn alloys [12,72]. Superimposed on Fig. 11 are two lines representing the experimental creep data for the Al 7005-

LI and LANGDON:

102 _

1 I I11111,

I

CREEP

BEHAVIOR

I I I I1111,

OF 7005 ALUMINUM

I I

ALLOY

I ! ,,!,,I,

1151

I , ,,I”U

Al 70052Ovol%~O,(p)

Al 606 I-2Ovol%A&O,(p)

4

Fig. 11. The transition between class M (climb control) and class A (glide control) for solid solution alloys: superimposed on the diagram are two lines representing the data obtained for the Al 700520vol.%A1203(p)

composite

in the present investigation and 20vol.%Al~03(p) composite

20vol.%A1,03(p) composite used in the present investigation and the Al 6061-20~01. %AlzOX(p) composite used earlier [l]. It is apparent that these two composites lie unambiguously within the regions of the diagram appropriate to the experimental results, with the composite having the Al-7005 matrix lying within the climb-controlled region with II close to 5 and the composite having the Al-6061 matrix lying within the glide-controlled region with n close to 3. Figure 11 therefore provides strong support for the proposal that the creep of metal matrix composites divides into two separate classes, class M and class A, with the measured behavior reflecting the dominant creep mechanism within the matrix alloy. 5. SUMMARY

data [l].

obtained

earlier

for an Al 6061-

inforced with 20 vol.% of irregularly shaped AlzOX particulates, where the value of n was close to 3 and the behavior was interpreted in terms of control by a dislocation glide process and class A behavior in the Al-6061 matrix alloy. 5. It is shown by analysis that both sets of data are consistent with the predicted transition from class M to class A behavior introduced earlier for the creep of unreinforced solid solution alloys. The results demonstrate that the creep of metal matrix composites is controlled by deformation of the matrix alloys and, as in solid solution alloys, there is a division into class M and class A types of behavior.

AND CONCLUSIONS

Creep tests were conducted on an Al-7005 metal matrix composite reinforced with 20 vol.% of irregularly shaped A&O3 particulates over a temperature range from 573 to 773 K. The experimental results show significant curvature in plots of the minimum creep rate vs stress, leading to high values for the apparent stress exponent and the apparent activation energy for creep. By introducing a threshold stress into the analysis, it is demonstrated that the true stress exponent, n, is close to 4.4 and the results are consistent with creep controlled by a dislocation climb process and class M behavior in the Al7005 matrix alloy. The results are compared with earlier creep data for an Al-6061 metal matrix composite, also re-

Acknowledgement-This work was supported by the U.S. Army Research Office under Grant DAAH04-96-l-0332.

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Metall.

Sci. Eng, 1997, Mater.

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F. A. and Langdon, T. G., 1495. and Karashima, S., Metal Sci.,

LI and LANGDON:

1152

CREEP BEHAVIOR OF 7005 ALUMINUM

8. Yavari, P. and Langdon, T. G., Acta metali., 1982, 30, 2181. 9. Endo, T., Nomura, T., Enjyo, T. and Adachi, M., J. Japan Inst. Metals,

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1967, 60, 51.

14. Murty, K. L., Mohamed, F. A. and Dorn, J. E., Acta metall., 1972, 20, 1009.

15. Park, K.-T., Lavernia, E. J. and Mohamed, Acta metall. mater.,

F. A.,

1990, 38, 2149.

16. Mohamed,

F. A., Park, K.-T. and Lavernia, E. J., Mater. Sci. Eng., 1992, A150, 21. 17. Pandey, A. B., Mishra, R. S. and Mahajan, Y. R., Acta metall. mater.,

1992, 40, 2045.

18. Kim, H. Y. and Hong, S. H., Scripta metall. mater., 1994, 30, 297.

19. Cadek, J., Oikawa, H., Sustek, V. and Pahutova, M., High Temp. Mater.

Proc., 1994, 13, 327. F. A., Metall. Mater. Trans., 1995, 26A, 3119. 21. Cadek, J., Oikawa, H. and Sustek, V., Mater. Sci. Eng., 1995, A190, 9.

20. Park, K.-T. and Mohamed,

22. Pandey, A. B., Mishra, R. S. and Mahajan, Y. R., Metall. Mater.

Trans.,

1996, 2lA,

305.

23. Zong, B. Y. and Derby, B., Acta Mater., 1997, 45, 41. 24. Mohamed, F. A. and Langdon, T. G., MetaN. Trans., 1974, 5, 2339.

25. Lagneborg, R. and Bergman, B., Metal Sci., 1976, 10, 20. 26. Bird, J. E., Mukherjee, A. K. and Dorn, J. E., in Quantitative Relation Between Properties and Microstructure, ed. D. G. Brandon and A. Rosen. Israel Universities Press, Jerusalem, 1969, p. 255. 27. Sherby, 0. D., Klundt, R. H. and Miller, A. K., Metall.

Trans., 1977, 8A, 843.

28. Gonzalez-Doncel, mater.,

G. and Sherby, 0. D., Acta metall.

1993, 41, 2797.

29. Mishra, R. S. and Pandey, A. B., MetaN. Trans., 1990, 21A, 2089. 30. Zhu, S. J., Peng, L. M., Ma, Z. Y., Bi, J., Wang, F. G. and Wang, Z. G., Mater. Sci. Eng., 1996, A215,

120. 31. Park, K.-T., Lavernia, E. J. and Mohamed, Acta metall. mater.,

F. A.,

1994, 42, 667.

32. Mishra, R. S., Pandey, A. B. and Mukherjee, A. K., Mater. Sci. Eng., 1995, A201, 205.

33. Cadek, J., Sustek, V., Kloc, L. and Evangelista, E., Mater. Sci. Eng., 1996, A215, 73. 34. Kloc, L., Spigarelli, S., Cerri, E., Evangelista, E. and Langdon, T. G., Metall. Mater. Trans., 1996, 27A, 3871. 35. Kloc, L., Spigarelli, S., Cerri, E., Evangelista, E. and Langdon, T. G., Acta mater., 1997, 45, 529. 36. Howson. T. E.. Mervvn. D. A. and Tien. J. K.. Metall. Trans., 1$80, llA, 1599.

37. Howson, Metall.

T. E., Mervyn, D. A. and Tien, J. K., Trans., 1980, llA, 1609.

38. Clauer, A. H. and Hansen, N., Acta metall., 1984, 32, 269. 39. Kuchafova, K., Orlova, A., Oikawa, H. and Cadek, J., Mater. Sci. Eng., 1988, AlO2, 201.

40. Yeh, Y. H., Nakashima, H., Kurishita, H. and Yoshinaga, H., Mater. Trans. JIM, 1990, 31, 284. 41. Mishra, R. S., Paradkar, A. G. and Rao, K. N., Acta Metall. Mater.,

1993, 41, 2243.

ALLOY

42. Fuentes-Samaniego, R., Nix, W. D. and Pound, G. M., Phil. Mug., 1980, 42, 591. R. and Nix, W. D., Scripta 43. Fuentes-Samaniego, Metall., 1981, 15, 15. 44. Herring, C., J. Appl. Phys., 1950, 21, 437. 45. Weertman, J., Trans. Am. Sot. Metals, 1968, 61, 681. 46. Darken, L. S., Trans. AIME, 1948, 175, 184. 47. Stoebe, T. G., Gulliver, R. D., Ogurtani, T. 0. and Huggins, R. A., Acta metall., 1965, 13, 701. 48. Hilliard, J. E., Averbach, B. L. and Cohen, M., Acta metall., 1959, 7, 86. 49. Li, Y., Nutt, S. R. and Mohamed, F. A., Acta mater., 1997, 45, 2607. 50. Friedel, J., Dislocations Pergamon, Oxford. 1967, p,

356. 51. Liiffler, H., Kovacs, I. and Lendvai, J., J. Mater. Sci., 1983, 18, 2215. 52. Li, Q., Johnson, E., Johansen, A. and SarholtKristensen, L., J. Mater. Sci., 1993, 28, 691. 53. Bartges, C. W., Scripta metall. mater., 1993, 28, 1039. 54. Chou, M.-C. and Chao, C.-G., Metall. Mater. Trans., 1996, 27A, 2005. 55. Ferragut, R., Somoza, A. and Dupasquier, A., J. Phys.: Condens. Matter, 1996, 8, 8945. 56. Hong, S. I. and Gray, G. T., Acta metall. mater., 1992, 40, 3299. 57. Orowan, E., in Dislocations in Metals, ed. M. Cohen.

AIME, New York, NY, 1954, p. 131. 58. Kocks, U. F., Phil. Mug., 1966, 13, 541. 59. Shewfelt, R. S. W. and Brown, L. M., Phil. Mug., 1977, 35, 945. 60. Arzt, E. and Ashby, M. F., Scripta metall., 1982, 16, 1285. 61. Arzt, E. and Wilkinson, D. S., Acta metall., 1986, 34, 1893. 62. Arzt, E. and Riisler, J., Acta metall., 1988, 36, 1053. 63. Rosier, J. and Arzt, E., Acta metall. mater., 1990, 38, 671. 64. Mishra, R. S., Nandy, T. K. and Greenwood, G. W., Phil. Mug. A, 1994, 69, 1097. 65. Weertman, J., J. Appl. Phys., 1957, 28, 1185. 66. Weertman, J., Trans. AZME, 1960, 218, 207. 67. Cottrell, A. H. and Jaswon, M. A., Proc. R. SOL,

1949, A199, 104. 68. Fisher, J. C., Acta metall., 1954, 2, 9. 69. Suzuki, H., Sci. Rprts Res. Inst. Tohoku Univ., 1952, A4, 455; ibid., 1955, A7, 194. 70. Snoek, J. L., Physica, 1942, 9, 862. 71. Schoeck, G., Phys. Rev., 1956, 102, 1458. 72. Mohamed, F. A., Mater. Sci. Eng., 1983, 61, 149. 73. Sellars, C. M. and Quarrell, A. G., J. Inst. Metals, 1961-1962,90, 329. 74. Cohen, J. B. and Fine, M. E., J. Phys. Radium, 1962, 23, 749. 75. King, H. W., J. Mater. Sci., 1966, 1, 79. 76. Cowley, J. M., Phys. Rev., 1950, 77, 669. 77. Rudman, P. S. and Averbach, B. L., Acta metall., 1954, 2, 576. 78. Hilliard, J. E., Averbach, B. L. and Cohen, M., Acta metall., 1954, 2, 621. 79. Smallman, R. F. and Dobson, P. S., Metall. Trans., 1970, 1, 2382. 80. Gallagher, P. C. J., Metall. Trans., 1970, 1, 2429. 81. Kannan, V. C. and Thomas, G., J. Appl. Phys., 1966, 37, 2363. 82. Lee, E. U., Kranzlein, H. H. and Underwood, E. E., Mater. Sci. Eng., 1971, 8, 336.

CREEP

LI and LANGDON: APPENDIX Estimation

BEHAVIOR

A

of the DifSusion Coeficients

ALLOY

APPENDIX for Climb and Glide

In an early analysis of creep in binary solid solution alloys [3], the diffusion coefficient for class M creep behavior and control by dislocation climb, D,, was taken as the Herring-Weertman weighted diffusion coefficient, DHW, given by [44,45]

An Examination

1153

B

of Possible Mechanisms the Threshold Stress

Responsible for

Several mechanisms have been proposed to explain the origin and magnitude of the threshold stresses in materials where mobile dislocations pass through arrays of particles. These mechanisms may be summarized briefly as under:

D;D;

D,=DHW=

l-41)

I"@'AD; + XBD~)

(1) When dislocations bow out between particles, the threshold stress is equal to the Orowan stress, rorw, which is given by [57.58]

where Da and 0; are the tracer diffusivities for the A and B atoms in the AB alloy, XA and Xn are the atomic fractions of A and B in the alloy andfis a correlation factor having a value of -0.78 for f.c.c. crystals. In class A creep behavior and control by dislocation glide, where creep takes place through some form of viscous drag process such as the dragging of solute atom atmospheres [2,3], the diffusion coefficient for glide, D, was taken as the-Darken chemical interdiffusivity for the solute atom, D. given by 1461

D~=D=jX;O;,+X,D;i(i+~)

Gh ~orw = 0.84 (). _ do) where A is the interparticle particle size.

spacing

and dp is the average

(2) When a back stress, rb, is needed to create an additional segment of dislocation as it surmounts the obstacle by local climb, the threshold stress is given by ]59, '501

(A2)

where FA is the activity coefficient for the A atoms. More recently, Fuentes-Samaniego et al. [42,43] re-examined the diffusion coefficients for solid solution alloys and derived diffusion coefficients for creep controlled by climb, Dz, and creep controlled by glide, 0;. Their relationships are given by

Df

OF 7005 ALUMINUM

=~(.X’d$ +XBD~)

(3) When there is a detachment stress, zd, associated with detaching a dislocation from an attractive particle, the threshold stress is given by [61-63]

where K, is a relaxation the attractive interaction location.

parameter which depends upon between the particle and the dis

643) (4) When there is a dissociation stress, f&s, due to the dissociation oflattice dislocations into interfacial dislocations when they enter the matrix-particle interface and surmount the obstacle by climb, the threshold stress is given

and

by F41 D; = (

x,$!;ADi)

644)

(’ + s)

(B4)

Equations (A3) and (A4) have been used successfully to analyze the experimental creep behavior for an Al-2.2 at.% Zn alloy [ll] and an Al-10 at.% Zn alloy [12]. For the Al 7005-20vol.%A120s(p) composite, the values of D,* and 0: were estimated by assuming an average concentration of Zn in the Al-7005 matrix alloy of -4.5 wt% from Table 1, equivalent to a Zn concentration, c, of 0.021, and using diffusion data for Al [47] and Zn [48] in Al-Zn ajloys to *calculate the relevant diffusion coefficients, DA, and DZn, listed in Table Al for the four different testing temperatures. These values of Da, and Di: were then used in equations (A3) and (A4) to estimate D, and Oi, and Table Al lists the calculated values together with the individual values of the ratio Of/D: for each testing temperature.

Table Al. Calculated diffusion coefficients for climb and glide, D’, and D*,, in the Al-7005 matrix alloy I

*

DA1 T

(K) (m* s-‘)

573 623 673 773

3.1 2.6 1.6 2.9

x x x x

10-l’ lo-l6 lo-l5 1O-‘4

1

DZ” (In2 s-1)

DC (m* s-l)

5.0 3.3 1.9 2.9

5.3 4.1 2.6 4.4

x x x x

IO-l6 1O-‘5 lo-l4 IO@’

x x x x

10m” lo-l6 lo-l5 lo-l4

0:/D; 3.4 x 2.4 x 1.4x 2.2 x

IO-l6 lo-” lo-l4 1Om’3

0.16 0.17 0.19 0.20

where rp is the particle

radius.

In order to make use of these models, it is necessary to estimate the appropriate values for dp and i. Inspection by TEM revealed precipitates having sizes in the range of -15-lOOnm, as documented in Section 3.3, but most of the particles were small and it is therefore reasonable to take d,r20 nm. The interparticle spacing is given by [35] 1 i” = ~ 2(Npdp)‘/’ where Np is the precipitate

density

defined as

where rzr is the measured number of particles within an area A, and t’is the foil thickness. Measurements by TEM gave N,r500~~m-~ so that 1~‘ 160 nm. Therefore, substituting these values for I and dp (=2r,) into equations (Bl)-(B4) with K,=O.85 [31,49], taking b = 2.86 x 10-‘” m for pure Al and using the shear modulus for pure Al (in MPa) defined as [6] G = (3.022 x 104) - 16T,

(B7)

1154

LI and LANGDON:

CREEP BEHAVIOR OF 7005 ALUMINUM

the normalized values for the theoretical threshold stresses are listed in Table Bl for a testing temperature of 623 K. These theoretical values lie within the range of -1.0 x lo4 to -1.7 x 10m3 whereas the experimental value for the normalized threshold stress at this temperature is r,/G = 2.5 x 104. Thus, the magnitude of the experimental threshold stress is consistent with the expectations from the theoretical models and, as in experiments on a PM 2124 Al alloy [49], the measured value of the normalized threshold stress lies within the range of r&G < r,/G < r,,/G. Table Bl. Estimated values for the threshold stress at a temperature of 623 K Mechanism

Normalized threshold stress

Orowan stress

zo,/G

Back stress due to local climb Detachment stress Dissociation stress

z~/G = 5.4 x lo4 Q/G = 9.5 x lOA Q>~/G= 1.0x lOA

= 1.7 x 1O-3

ALLOY

(iv) The stress-induced local ordering of solute atoms through the Snoek interaction [70,71] where A” = A”%I> and (v) The antiphase boundary alloys [66] where A” = A:,,.

interaction

Considering these five processes, the latter applies only in ordered alloys and therefore can be neglected and Mohamed [72] showed, through a detailed analysis of a large number of alloys, that the Snoek interaction makes a very minor contribution in most systems including in AlZn and Al-Mg alloys. Therefore, it is necessary to examine only the first three of these five processes. From the theoretical descriptions of these processes, the values of the constant A” in equation (C2) are given by [67,68,73,74] e2cb5G2 A% =w s

APPENDIX C

A,,

Am Estimation of the Criterionfor the Class M-class A Transition in Metal Matrix Composites In the creep of solid solution alloys, it is well established that creep by viscous glide in class A occurs under a condition which is given by [3]

$&3)2< qg3(2)($’

=

B

G(i)’

b2

and

2: g,

A,,

v(rB

S

-

rA)2

(i) the segregation of solute atoms to moving dislocations through the Cottrell-Jaswon interaction [67] where A” = A” CJI

(ii) The destruction of short-range ordering through the Fisher interaction [68] where A” = A$, (iii) The chemical interaction of solute atoms with dissociated dislocations through the Suzuki interaction [69] where A” = A” S,

(C5)

p4

with -8AH,/2

(W

where x is the degree of local order, y is the bond energy, V is the molar volume, rA and ra are the stacking fault energies for the solute and the solvent, respectively, A S/ AS,, is the deviation of the entropy of mixing from the ideal value, AH,,, is the entropy for mixing of the 50% alloy, B is a constant having a value of -14 for f.c.c. crystals and B” is a constant having a value of -0.9. When several viscous drag processes operate, each mechanism has a different value for A”; and these processes operate sequentially so that the slowest process provides the major dragging force [72]. Therefore, the constant A” in equation (C2) may be replaced by the summation of the individual values of A” so that the shear strain rate is expressed as

03

+j,=@‘G;

where B is a constant having a value of -0.47, A” is a constant determined by the relative contributions from the different viscous glide processes, and the shear strain rate and shear stress are related to the normal strain rate, B, and the normal stress, u, through the expressions 9 = 38/2 and t = (r/2. Following the approach of Weertman [66], it is necessary to consider the following five different viscous drag processes:

(C3)

XAXEJXY

F

(Cl)

where e is the solute-solvent size difference, c is the concentration of the solute, r is the stacking fault energy of the alloy, B is a constant having a value which was estimated in the earlier analysis by using creep data for an Al-3% Mg alloy [3] and CIis a constant which incorporates the relative contributions from the different possible viscous glide processes. In order to determine the value of c(for use in this analysis, it is necessary to examine the nature of the different viscous drag mechanisms. When creep occurs under conditions where dislocation glide is rate-controlling, the shear strain rate, js, is given by F51

?g=

in ordered

p=

l

A&+A;+A;

G3)

For Al-Mg alloys, A% = 0 and it has been shown that A& > A: [72]. Therefore, fi = l/A& and, consistent with the earlier analysis for Al-Mg solid solution alloys [3], c( = 1 in equation (Cl). For Al-Zn alloys, however, analyses show that the Fisher and Suzuki interactions make a significant contribution to the viscous drag process in addition to the Cottrell-Jaswon interaction [12,72] and it is necessary therefore to evaluate the individual values for A”CJ> A”F and A” Taking c = sd.021 for the Al-7005 matrix alloy with e = 0.02 [75], and putting D, equal to the values of 0; tabulated in Table Al, it is possible to use equation (C3) to calculate values of A& for each testing temperature and these values are listed in Table Cl. The importance of the Fisher interaction may be estimated by determining y from the following expression [76]:

CREEP

LI and LANGDON:

c-1 %

ex’ kT

_

(XA

+ xB,d(xB xAxB(l

+ XAX) -

BEHAVIOR

(0)

Xj2

Taking ,y = 0.082 [77] and again using 0; from Table Al, values of A; were calculated from equation (C4) and listed in Table Cl. The Suzuki interaction may be estimated from equations (C5) and (C6) by taking A S/A&-1.2 [78], Vs 10 cm3 mol-‘, (aJ’/&) v‘ - 0.78 cm mol-’ for dilute Al-Zn alloys, AHlIZ = 3.36 kJ mol-’ [77,78], rA= 140 mJ mm2 [79] and Fa = 200 mJ m-* [80] to give the values of A: listed in Table Cl. The ratios of Ag/A& and Ag/A& are also shown in Table Cl and they provide information on the relative importance of each of these viscous drag processes. Finally, these values may be used to estimate / through equation (C8) and hence to determine the appropriate values of r in equation (Cl) for the Al-7005 matrix alloy at each testing temperature; these values are recorded in the last column of Table Cl and thev confirm that the Fisher and Suzuki interactions make a ‘non-trivial contribution in the Al-Zn matrix alloy. A criterion for a transition between class M and class A behavior may be developed for metal matrix composites bv noting that, as outlined in Section 4.1. it is necessary to replace the shear stress, r, in equation (Cl) with the e&ective shear stress, (t - 7,). Thus, the transition occurs at the condition given by

I’=

~(gJ3(g&y2

(C10)

OF 7005 ALUMINUM

ALLOY

1155

The present analysis is concerned with AI-6061 and Al7005 matrix alloys where Mg and Zn are the respective major alloying elements. Equation (ClO) can be used for the Al-6061 matrix alloy by taking, as in the earlier anaiysis for_ an Al-3% Mg solid solution alloy [3], D,=DHW, with e = 0.1208 [75], F = 190 D,=D and D,/D,=l, mJ rnd2 [81] and c= 0.011 as a representative value for this alloy. Using the same approach for the Al-7005 matrix alloy, consistency with the creep analysis requires in Table Al, that D, = D,* and D, and 0; as calculated with e = 0.0195 [75], Tr 170 mJm-* estimated by interpolation between I’a and the experimental value of r = 54 mJ m-’ reported for an Al-9.9 at.% Zn alloy [82], and again using c = 0.021 for a Zn concentration of -4.5 wt% in this alloy. The validity of equation (ClO) is demonstrated in Section 4.3 using a graphical presentation.

Table Cl. Calculated matrix alloy

573 623 673 773

2.6 x 3.2x 4.7 x 2.2 x

lO-3 lOA 10m’ IOP

8.4 1.3 2.4 1.7

x x x x

values of A” and G(for the Al-7005

10m4 lOA lO-5 1oP

3.4 x 4.5x 7.1 x 3.9 x

IO+ lO-5 lO-6 lo-’

0.32 0.41 0.51 0.80

0.13 0.14 0.15 0.18

0.69 0.65 0.60 0.50