Electron microscopy observations of deformation twinning in a precipitation hardened copper-titanium alloy

ScriptaMaterialis, Vol. 35, No. 12, pp. 1403-1409,1996 Elscvier Science Ltd Copyright0 1996 Acts MetallurgicaInc. printed in the USA. All rights reserved 1359-6462196$12.00 + .OO

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ELECTRON MICROSCOPY OBSERVATIONS OF DEFORMATION TWINNING IN A PRECIPITATION HARDENED COPPER-TITANIUM ALLOY T. Radetic, V. Radmilovic* and W.A. Soffa Department of Materials Science and Engineering University of Pittsburgh, Pittsburgh, PA 1526 1 (Received February 13,1996) (Accepted July 1, 1996) Introduction Copper-titanium alloys in the range 1 to 5 wt% titanium can exhibit properties comparable to the wellknown Cu-Be alloy series after suitable heat treatment (1). Age hardened Cu-Ti alloys represent a potentially important substitute for the conventional high-strength Cu-Be materials in a variety of applications, particularly in the electronics industry. Processing of the traditional Cu-Be materials involves many .problems primarily associated with the toxicity of Be. From the early 1930’s until present, a burst of research has been periodically directed at the Cu-Ti alloys with much of this work focused primarily on the complex decomposition phenomena exhibited by supersaturated coppertitanium solid solutions (2,3,4). It is now well-established that alloys in the range 2-5 wt% Ti decompose by a spinodal mechanism at large supersaturations accompanied by ordering in the T&rich regions (4). The resultant metastable, coherent Cu4Ti precipitate phase (p’) has the Dl, superstructure (2, 3). The precipitation reaction gives rise to a marked age hardening and the dispersion of p’ particles can enhance the yield strength up to several times that of pure copper (Cu) with a similar grain size. In spite of the rather extraordinary age hardening response of the Cu-Ti alloys there has been a striking paucity of systematic work reported on the relationship of the microstructure and mechanical properties. In fact, the strengthening mechanism responsible for the high strengths developed in these alloys is yet to be elucidated. The flow and fracture behavior of these particle-hardened alloys is particularly interesting in that both single crystals and polycrystals show an unusual propensity to twin profusely after small amounts of plastic flow by slip (5, 6). Recently, suitably oriented Cu-Ti-Al (-lwt?!! Al; the Al addition lowers the stacking fault energy) single crystals containing coherent Cu4Ti (Dla) particles have been reported to yield by twinning at the onset of plastic flow and it has been concluded that the ordered phase facilitates deformation twinning (5). In this paper the preliminary results of conventional electron microscopy (CTEM) and highresolution electron microscopy (HREM) studies of the fine-scale structure of the mechanical twins which form in particle hardened copper-titanium alloys are reported. *Department of Phy:rieal Metallurgy, Karnegijeva 4, P.O. Box 494, University of Belgrade, 11001 Belgrade, Yugoslavia

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Experimental Tensile specimens were prepared from a cold rolled Cu-4.2 wt% Ti alloy sheet material with gage length dimensions 25.4 x 6.4 x 0.2 mm. The alloy specimens were solutionized at 875°C for two hours and then quenched into ice brine. They were then aged at 375°C for times ranging from 30-1500 mm. All heat treatments employed encapsulating the specimens in quartz tubes under an argon atmosphere. The grain size of the heat treated alloys was approximately 60-70 pm. Tensile testing was performed on a hydraulically driven machine at a strain rate i=6.7 x lOA se’.For each aging time specimens were deformed plastically - 5% or until fracture. Discs (3 mm dia.) were punched from the gage length of the tensile specimens and mechanically polished to approximately 0.08 mm and then electropolished for the TEM studies using a solution of 75 g Cr03, 375 ml of glacial acetic acid and 20 ml of HzO. The optimum polishing condition was approximately 50 mA and 20-30 V at room temperature. Thin foils were examined in a JEOL 200 CX transmission electron microscope at 200 kV and a JEOL ARM-100 microscope at 800 kV. Results and Discussion A maximum yield strength of approximately 680 MPa was developed in the Cu-4.2 wt% Ti alloy aged for the 1500 min at 375°C. The microstructure and deformation substructure of these samples were characterized. Examination of the microstructure of the alloy aged for the 1500 min at 375°C by conventional TEM (BF and SAD), reveals the presence of p’ particles within the Cu-rich matrix. The diffraction pattern shows superlattice spots of the ordered D 1, phase (p’). As a consequence of the large extinction distance of the 0’ particles, 5 is approximately 350 nm (2), superlattice reflections can be seen only in overexposed diffraction patterns. Their low intensity makes dark-field (DF) imaging difficult because of the long exposure times required except for very long aging times. The coherent p’ particles at peak strength are essentially prolate spheroids with c=35 nm and az17.5 nm. The deformation substructure of the samples deformed 5% and deformed to fracture (15% deformation) shows that profuse twinning occurs during the deformation process. Twinning is not a common deformation mode for two-phase f.c.c.-based alloys with coherent, ordered particles, since the mechanism that produces twins in the matrix generally does not produce twins in the ordered structure. In the case of the j3’particles, due to the crystallography of the superstructure (Dl,) and its orientation relationship with the f.c.c. matrix, the twinning mechanism in the matrix can produce true twinning in the precipitate preserving the superlattice (7). Multiple twinning is generally observed within the grains as shown in Figure 1. The twins are often quite long, extending across the grains. When a twin/twin intersection occurs the intersecting twins generally continue to propagate, but sometimes an incident twin terminates on a obstacle twin. Conventional transmission electron microscopy reveals that the thickness of deformation twins is of the order of 10-2000 nm, which is in agreement with the results of another authors (5, 6). One of the mechanisms of twin thickening is the coalescence of thinner twins (Figure 2). Streaking along ~11 l> traces in the diffraction patterns indicates the presence of very fine tWillS.

High resolution electron microscopy shows the presence of very fine twinned regions whose thickness is of the order of l-3 mn. At the same time, HREM indicates that the twins observed by conventional TEM are actually bundles of very fine twins. HREM does not reveal any structural perturbation within the deformation twins which could indicate the presence of the ordered precipitate. This is not surprising considering the large extinction distance of the Dl, phase. Results of the HREM study show that the elementary process of deformation twin growth involves the G-ledge or glide mechanism. Applying the Frank circuit approach (8, 9), the Burgers vectors of dislocations associated with the ledges, the so-called twinning dislocations, were identified and they were of the same type as the

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Figure 1. Transmission electron micrograph of deformation twinning on multiple systems. Bright field image; [l lo],.

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Figure 2. Transmission electron micrograph of coalescence of deformation twins. Bright field image; [ 1lo],.

Shockley parti,al dislocation, bt=1/6<1 12>. ( The notation S and F refer to the conventional start and finish of the Frank circuit analogous to a Burgers circuit.) The ledge thickness varies between one to a few atomic layers. In Figure 3, the microtwin growth associated with a one atomic layer thick G-ledge can be seen. Along the twin, the thickness varies from the six to nine atomic layers. The value of the Burgers vector

Figure 3. HREM micrograph of a deformation twin; [ 1lo],.

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Figure 4. HREM micrograph of a deformation twin. Different types of ledges and dislocations associated with them are revealed; r1101za.

of the twinning dislocation is determined to be l/12-4 12>, revealing that it is associated with the edge component of the so-called 30’ disiocation (l/64 12> dislocation whose Burgers vector is 30” inclined to the electron beam, so it does not lie in the projection plane). The ledges and associated twinning

Figure 5. HREM micrograph of a twin/matrix interface. A ledge, two atomic layers thick is associated with the interface; [I lO]za.

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Figure 6. HREM micrograph of a twin/matrix interface. A pure ledge, three atomic layers thick is associated with the interface; [l IOJza. dislocations (a) and (b) are on the same twin/matrix interface and the ledge and associated twinning dislocation (c) #areon the opposite interface. Under the action of a positive shear stress the motion of (a) and (b) will cause twin growth and (c) will act as a detwinning dislocation; if the sign of the stress is reversed the motion of (a) and (b) will cause detwinning and (c) will cause twin growth. The results of HREM show that different types of steps (9) are associated with twin interfaces. Figure 4 reveals the presence of a few ledges within a short distance. The dislocation associated with the G-ledge (a) corresponds to the twinning dislocation whose Burgers vector lies in the (110) plane of the micrograph, i.e. it is a 90” dislocation. The step (b) on the opposite twin/matrix interface is also a G-ledge, but the dislocation associated with it is characterized as a 30” dislocation, showing that twinning dislocations with different Burgers vectors are associated with the same twin. This implies that either dislocation (a) or (b) is a detwinning dislocation. The nature of ledge (c) indicates that (c) and (b) can be formed by dissociation of a unit dislocation on the twin/matrix interface and dislocation (b) could have a detwinning character. Ledge (c) is associated with a dislocation with bF1/3<111> that is normal to the t\vin boundary. For the motion of this ledge along the {11 l} glide planes climb is necessary, so it is called a C-ledge. Figure 5 shows a two atomic layer thick G-ledge. Two dislocations are essentially associated with this particular ledge on different levels. The components of the Burgers vectors in the plane of projection are each 1/12<112> indicating edge components of a 30” dislocation. An example of a pure ledge (a pure ledge essentially has no dislocation character) can be seen in Figure 6. It is tlhree atomic layers thick. The Frank circuit associated with this ledge is closed as in the case of perfect crystal, indicating the net Burgers vector is zero (characteristic of a pure ledge). A mechanism of Formation of this type of step was first proposed for b.c.c. crystals by Sleeswyk (lo), and later for the f.c.c. structure by Hirth and Balluffl(9). The pure ledge can be formed by emissary slip, i.e. dissociation of a three layer thick G-ledge possessing a Burgers vector 3bt, into a lattice dislocation and a pure ledge. The lattice dislocation can slip away, leaving behind a stress free ledge. An alternative, but less probable mechanism is the cancellation of the G-ledge by a dislocation with opposite Burgers vector, -3 bt (9).

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Figure 7. HREM micrograph of a complex dislocation confguration close to a twin/matrix interface; [ lOi],.

The fact that the deformation twins are distributed uniformly over the grains and since little dislocation substructure could be observed by conventional TEM, suggests that twinning is a primary mode of deformation in the sample aged 1500 min. It does not mean that dislocations, besides twinning dislocations of type 1/6<112>, are not present in the structure. Because of the significant strain fields around CtuTi precipitates caused by a mismatch between the Cu4Ti particles and matrix, as well as due to the large amount of twinning, the individual dislocations were difficult to resolve. However, HREM study reveals the presence of other dislocations. In Figure 7 there is an extended 60” dislocation in the (110) projection. It is dissociated into the 30” and 90” Shockley partial dislocations. The core of the 30” dislocation is localized (b), i.e. a perfect intrinsic stacking fault order is exhibited in the next atomic layer. On the other hand the 90” (a) dislocation is characterized by a delocalized core, so a perfect intrinsic stacking fault sequence is established within four atomic layers. Strongly localized 30” and delocalized 90” dislocations are common in high-resolution images of dissociated 60” dislocations (11). The width of the extended dislocation is 30 atomic layers. This separation cannot be used for the determination of the SFE due to the complex configuration; from the one side the extended dislocation is bounded by a twin boundary, on the other the 30” component (b) is locked by the interaction with a 60” unit dislocation (c), forming a 1/6<141> stair rod configuration (d). A total Burgers vector of this complex dislocation configuration is . Conclusions These electron microscopy observations of deformation twins in a particle-hardened Cu-4.2wl% Ti alloy showed that twin growth occurs via the motion of so-called G- ledges on the twin interface and via coalescence of fine twins. The G-ledges exhibit a Burgers vector bl=a/6<1 12>. Dislocations with an edge character (90”) and mixed character (30’) have been observed. In addition to the G-ledges, so-called C-ledges and pure ledges were observed.

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The Cu4Ti (p’) particles were not imaged in HREM because of the large extinction distance (4 characterizing the D 1Bordered phase.

Acknowledgements The authors would like to acknowledge the assistance of the staff at the NATIONAL CENTER FOR ELECTRON MICROSCOPY at Berkeley in carrying out the HREM experiments. Also, the technical assistance and advice of Mr. Tom Nuhfer at the Department of Materials Science at CMU is gratefully acknowledged. Finally, the authors would like to thank Professor R. Gronsky of the Department of Materials Science at Berkeley for his helpful discussion of the electron microscopy results. This work was supported in part by the National Science Foundation (NSF), Division of Material Research (DMR). References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11.

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