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ISSN : 1225-0562(Print)
ISSN : 2287-7258(Online)
Korean Journal of Materials Research Vol.29 No.6 pp.343-348
DOI : https://doi.org/10.3740/MRSK.2019.29.6.343

Dynamic Precipitation and Substructure Stablility of Cu Alloy during High Temperature Deformation

Chang-Suk Han†, Dong-Nyeok Choi, Sung-Yooun Jin
Dept. of ICT Automotive Engineering, Hoseo University, 201, Sandan7-ro, Seongmun-myeon, Dangjin City, Chungnam 31702, Republic of Korea
Corresponding author E-Mail : hancs@hoseo.edu (C.-S. Han, Hoseo Univ.)
April 1, 2019 June 14, 2019 June 14, 2019

Abstract


Structural and mechanical effects of the dynamical precipitation in two copper-base alloys have been investigated over a wide range of deformation temperatures. Basing upon the information gained during the experiment, also some general conclusion may be formulated. A one concerns the nature of dynamic precipitation(DP). Under this term it is commonly understood decomposition of a supersaturated solid solution during plastic straining. The process may, however, proceed in two different ways. It may be a homogeneous one from the point of view of distribution and morphological aspect of particles or it may lead to substantial difference in shape, size and particles distribution. The effect is controlled by the mode of deformation. Hence it seems to be reasonable to distinguish DP during homogeneous deformation from that which takes place in heterogeneously deformed alloy. In the first case the process can be analyzed solely in terms of particle-dislocation-particle interrelation. Much more complex problem we are facing in heterogeneously deforming alloy. Deformation bands and specific arrangement of dislocations in form of pile-ups at grain boundaries generate additional driving force and additional nucleation sites for precipitation. Along with heterogeneous precipitation, there is a homogeneous precipitation in areas between bands of coarse slip which also deform but at much smaller rate. This form of decomposition is responsible for a specially high hardening rate during high temperature straining and for thermally stable product of the decomposition of alloy.



초록


    Hanseo University

    © Materials Research Society of Korea. All rights reserved.

    This is an Open-Access article distributed under the terms of the Creative Commons Attribution Non-Commercial License (http://creativecommons.org/licenses/by-nc/3.0) which permits unrestricted non-commercial use, distribution, and reproduction in any medium, provided the original work is properly cited.

    1. Introduction

    The strain hardening and precipitation hardening are the most efficient ways by which the strength of the alloy may be improved.1-7) Simultaneous use of both of these means is of particular interest because of multiplicity of variants of strain/aging conditions which can be used leading to different mechanical properties. An obvious feedback between the mode of plastic deformation, formation of the dislocation substructure and decomposition of the solid solution makes that the structure of an alloy which is formed during a complex thermo-mechanical treatment influences: 1) the mechanical behavior during a high temperature deformation of an unstable solid solution (dynamic precipitation: DP) and 2) the level and stability of the structure(properties) after a thermo-mechanical treatment.8-11) For some practical applications the temporal stability of the structure is not less important than the hardening level.

    In this work we concentrate upon the interrelation between the mechanical characteristics and the mode of a high temperature deformation of the alloys under a DP conditions and the structure and properties of post treatment metal. Variety of alloys which are suitable for the DP treatment make that the choice of the material to the experimental study is a particular problem, especially if one wants to come to some general conclusions. Hence in the experiment we use the copper base alloys which show similar properties(hardness) in a fully annealed state but differ significantly in the solid solution hardening level(supersaturated solid solution or as we refer to latter; solution treated alloy: ST) and in the rate and level of the precipitation hardening.

    The development and stability of a dislocation/precipitation substructure in the course of a high temperature deformation of an unstable solid solution and their effect on the mode of deformation is studied by comparison the mechanical characteristics of supersaturated and “over aged”(OV) alloys. The term “over aged” means the reference material of a stable solid solution and stable second phase particles structure in a high temperature deformation conditions (temperature and time of straining).

    The stability of structure generated during the dynamic precipitation(in terms of the hardness variation) is also compared with the stability of the structure formed by cold working of the supersaturated alloy and then statically aged.

    2. Experimental Procedure

    2.1 Material

    Composition of two commercial purity Cu-alloys used in experiments was given in Table 1. Hot rolled sheets of 13 mm in thickness were cut along rolling direction and square shaped rods 13×13 mm were then cold rolled on a shape mill to the diameter of 10 mm. Specimens for compression test 12 mm long with convex necks of 0.2 mm deep on both sides were machined from the rods and then annealed at the temperature given in Table 1. Solution treated alloy was used for experiments on dynamic precipitation during hot compression tests. Over aged alloys were used as the reference material at which the solid solution decomposition was strongly limited during hot compression tests. The structure stability tests were performed out by means of micro-hardness measurements on ST samples of B-alloy deformed at different temperature and strain rate given in Table 2.

    2.2. Testing

    Compression tests at the temperature range of 620- 1,170 K were performed out on a modified INSTRON machine with constant true strain rate of 1.3×10−3 s−1 and 1.3×10−2 s−1 (Fig. 1). The flaked graphite has been used to reduce the friction between a specimen and anvils. Specimens were quenched in water approx. 1 second after the deformation had stopped.

    The hardness of hot deformed specimens was measured at room temperature using Vickers micro-hardness tester and a indenter load of 200 g. Longitudinally sectioned specimens were electropolished before testing(Fig. 2).

    2.3 Structure observations

    Optical microscopy observations were performed out on specimens sectioned along compression axis. Electropolished specimens were chemically etched in a solution of l0 g FeCl3 in 30 ml HCl and 120 ml methanol. For electron microscopy observations thin slices of the material were cut of along specimen axis and thin foils were prepared using a standard technique. Some structures of hot deformed alloys were shown in Figs. 3-4.

    2.4 Annealing of B-alloy

    Hot deformed specimens of ST alloy B were annealed at constant temperature to compare their hardening or softening kinetics with these for cold deformed material. Initial deformation conditions for specimens used in the experiment are listed in Table 2. The kinetics of hardening (softening) was shown in the Fig. 5.

    3. Results and Discussion

    In Figs. 1(a) and 1(b) are shown some stress vs. strain curves recorded in compression at different temperature for alloy A and B respectively. Dashed lines represent characteristics typical for the OV alloy, while solid lines are the characteristics of the supersaturated solid solution (solution treated alloy-ST). Two types of the σ-ε characteristics may be distinguished. One, characterized of a steady flow stress behavior which sets up soon after beginning of strain at high temperature and the other, where after initially fast strain hardening the deformation becomes mechanically unstable(stress peak) and then the curves reveal very efficient global softening effect. This type of the behavior was observed only in solution treated alloy A.

    It is worth of noticing that the hardening effect(current flow stress) of this alloy at all temperatures of the test is higher than that in the OV alloy A. This is, however, not the case of the alloy B where the flow stress of the over aged alloy is higher(except of the very high temperature, above that used to overage the alloy) than in the material undergone dynamic precipitation.

    The hardness of the alloys, which were deformed up to 0.5 strain value under the conditions as in Fig. l(a) and Fig. 1(b) [Fig. 2(a) and Fig. 2(b) respectively] also varies in a similar way. Specifically in ST alloy A the highest hardness after high temperature deformation falls to the temperature range in which deformation is highly unstable. Almost constant value of the hardness after deformation at 750 K and 850 K [Fig. 2(a)] clearly suggests that the structure produced by a complex deformation-precipitation treatment is very much the same and is thermally stable.12) Its very high value indicates that no recrystallization was involved into structure stabilization. Decrease of the hardness after deformation at higher temperatures finds a justification in the proximity of the solvus line, what makes that the effect of precipitation hardening diminishes.

    The hardness of OV alloy A after similar deformation is never as high as in ST alloy and may be even 100 hardness units smaller after the deformation in the temperature range between 750-900 K. The hardness of the solution treated and hot deformed alloy B increases with temperature from about 90 to 110 HV units in the range of deformation temperature up to 850 K[Fig. 2(b)]. In turns, the hardness of the OV alloy is high (HV 120) and stable all over this range. Then, hardness of both ST and OV alloy decreases accordingly with temperature.

    The optical microscope observations of the structure of alloys after high temperature deformation provide first of all the information about the mode of the deformation at high temperature and in particular about the nature of the mechanical instability in solution treated alloy A. As may be seen from the Fig. 3(a) the instability of the plastic flow is caused by concentration of the deformation in shear bands. The Fig. 3(a) reveals shear bands in the sample of ST alloy A after ~42 % deformation at 820 K. Shear bands show a tendency to form in the close vicinity of a grain boundary from where they expand across several grains. The pattern of the structure obtained by etching indicates different rate of the decomposition of the alloy along grain boundaries and within gains.

    There were not macroscopic shear bands neither in OV alloy A nor in any variants of alloy B[Fig. 3(b)]. Optical observations also show a pancake shape of the grains after high temperature deformation, which points to the absence of the recrystallization, (T<1170 K).

    TEM observations show that heterogenization of the deformation in ST alloy A begins much before the peak stress is achieved. Already after 5% of deformation the channeling of the substructure composed of dislocations and very fine precipitates(basically not resolvable in the structure) by a coarse transsubstructural slip can be seen [Fig. 4(a)]. Basically, two different types of the substructure are observable. In the area of a macroscopic shear located along a grain boundary the structure consists of elongated precipitates, often arranged into two sets of parallel plates or rods. With ongoing deformation these particles become severely chopped as is shown in Fig. 4(a). Inside grains very messy substructure is observed, with a tendency toward formation of dislocation cells at higher temperature[Figs. 4(b), (c), (d)].

    The experiments show that a very important factor which controls the effect of DP is the mode of a plastic flow in an unstable solid solution. Among characteristic features of the alloy A there is a strong effect of a solution hardening(strong dislocation-solute atom interaction) and strong effect of the precipitation hardening(strong dislocationparticles interaction). Both, cause that a preferable form of the plastic flow is highly concentrated in slip bands [Fig. 3(a) and Fig. 4(a)]. Their extend is limited by the grain size, or else, that only efficient barriers for this form of slip are grain boundaries.

    The reason for the heterogenization of deformation is understood in terms of soft path for the dislocation motion. Such a soft path forms in result of shearing off solute enriched zones or/and coherent particles(typical for early stage of aging) what reduces their strength and leads to an avalanche-like movement of dislocations on softened this way plane(coarse slip phenomenon).13,14)

    Hence, precipitation to a large extent is pseudo-static process in the area between bands of concentrated slip. The growth of the particles accompanied by the lost of coherency may, however, eliminate this mode of slip giving rise to a very high rate of strain hardening.

    A parallel effect of initial heterogenization of flow in bands of coarse slip is different mechanism and rate of decomposition of a solid solution in the areas along grain boundaries and inside grains. The difference is seen in size distribution and morphological aspect of second phase particles. The pattern which is typical for a grain boundary region clearly shows a very fast growth of particles there, while their shape and arrangement may suggest a discontinuous precipitation. The presence of few sets of parallel particles arrangements in the same area[marked by arrows in the Fig. 4(a)] makes, however, such a conclusion highly superficial. One can rather incline toward the conclusion that the growth of particles is enhanced by pile-ups of a dislocations arrested by grain boundary. This suggestion receives some support from well known fact that redistribution of dislocations is strongly limited by presence of very fine particles. If so, than segregation of solute atoms in vicinity of a pile-ups appears pile-up stress relaxing mechanism. Another words, the presence of pile-ups at grain boundaries provides additional driving force for oriented grain nucleation and then enhanced growth of particles. The observation of preferred precipitation at shear bands in an aluminumbase alloy may be used here also to support this conclusion.

    The experiments also shows that very intense growth of particles in the area close to grain boundaries leads to the quick over aging(softening). This seems to explain why suddenly plastic flow concentrates along grain boundaries. The phenomena begin at very high stress and leads to catastrophic flow(transgranular shear banding) accompanied by mechanical unstability of material.

    The analysis of the mechanical performance and properties of OV alloy A and both variants(OV, ST) of alloy B leads to the other conclusion. Straining of an alloy having already stable particles always gives lower and less thermally stable properties(hardness) than equivalent straining of ST alloy. This means that concomitant generation of dislocation and particles produces very though and stable structure.

    The stability of the structure and properties of the alloy during post-deformation annealing depends upon two opposite effects: a) static recovery which reduces the hardness of the material and b) precipitation hardening which depends of the kinetics of a solid solution decomposition and distribution and particles volume fraction. Dynamic recovery processes operative during hot deformation at low strain rate rise stability of a dislocation substructure during annealing and reduce the effect of static recovery. The hardness-time curves shown in the Fig. 5 are in agreement with the effect mentioned above. The hardness of the hot deformed material is lower than for cold deformed material aged 6 min. at given temperature. However, after long aging(>100 min.) the hardness of cold deformed material falls bellow the value for hot deformed one.

    4. Conclusions

    Basing upon the information gained during the experiment, also some general conclusion may be formulated. A one concerns the nature of dynamic precipitation. Under this term it is commonly understood decomposition of a supersaturated solid solution during plastic straining. The process may, however, proceed in two different ways. It may be a homogeneous one from the point of view of distribution and morphological aspect of particles or it may lead to substantial difference in shape, size and particles distribution. The effect is controlled by the mode of deformation. Hence it seems to be reasonable to distinguish DP during homogeneous deformation from that which takes place in heterogeneously deformed alloy. In the first case the process can be analyzed solely in terms of particle-dislocation-particle interrelation. Much more complex problem we are facing in heterogeneously deforming alloy. Deformation bands and specific arrangement of dislocations in form of pileups at grain boundaries generate additional driving force and additional nucleation sites for precipitation. Along with heterogeneous precipitation, there is a homogeneous precipitation in areas between bands of coarse slip which also deform but at much smaller rate. This form of decomposition is responsible for a specially high hardening rate during high temperature straining and for thermally stable product of the decomposition of alloy.

    Acknowledgement

    This research was supported by the Academic Research fund of Hoseo University in 2018(20180366).

    Figure

    MRSK-29-6-343_F1.gif

    True stress vs. true strain curves received at true strain rate of 1.3×10−3 s−1 for (a) A-alloy and (b) B-alloy(solid line : solution treated alloy, dash line : over aged alloy). Deformation temperature was marked in the figure.

    MRSK-29-6-343_F2.gif

    Microhardness of solution treated alloys(ST: solid line) and over aged one(OV: dashed line) deformed with εt=0.5 versus temperature of deformation. (a) A-alloy, (b) B-alloy. The strain rate was marked in the figure.

    MRSK-29-6-343_F3.gif

    Structure of solution treated (a) A-alloy deformed at the temperature of 820 K with strain rate of 1.3×10−3 s−1 ; εt=0.51; (MSB; macro shear bands) and (b) B-alloy deformed at the temperature of 1070 K with strain rate of 1.3×10−3 s−1 ; εt=0.50.

    MRSK-29-6-343_F4.gif

    Solution treated (a) A-alloy deformed at the temperature of 820 K with strain rate of 1.3×10−3 s−1, εt=0.05(CSB: coarse shear bands, GB: grain boundary) and B-alloy deformed with strain rate of 1.3×10−3 s−1 at the temperature of (b) 970 K, (c) 770 K and (d) over aged alloy, deformed at 1170 K; εt=0.50.

    MRSK-29-6-343_F5.gif

    Microhardness of a B-alloy deformed approx. εt= 0.50 v.s. time of annealing at the temperature of (a) 710 K and (b) 870 K. Specimens were pre-deformed with strain rate and at temperature given in a Table 2.

    Table

    Composition and hot treatment of alloys (ST: solution treated alloy, OV: over aged alloy.)

    Deformation conditions for specimens of solution treated B-alloy used in annealing experiments.

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