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Microstructural evolution of nanostructured steel-based composite fabricated by accumulative roll bonding
Author’s Accepted Manuscript Microstructural evolution of nanostructured steelbased composite fabricated by accumulative roll bonding Roohollah Jamaati, Mohammad Reza Toroghinejad, Sajjad Amirkhanlou, Hossein Edris www.elsevier.com
PII: DOI: Reference:
S0921-5093(15)00547-X http://dx.doi.org/10.1016/j.msea.2015.05.025 MSA32353
To appear in: Materials Science & Engineering A Received date: 25 February 2015 Revised date: 7 May 2015 Accepted date: 8 May 2015 Cite this article as: Roohollah Jamaati, Mohammad Reza Toroghinejad, Sajjad Amirkhanlou and Hossein Edris, Microstructural evolution of nanostructured steel-based composite fabricated by accumulative roll bonding, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.05.025 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructural evolution of nanostructured steelbased composite fabricated by accumulative roll bonding
Roohollah Jamaati 1 Department of Mechanical Engineering, Babol Noshirvani University of Technology, Babol, Iran
Mohammad Reza Toroghinejad Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156–83111, Iran
Sajjad Amirkhanlou Young Researchers and Elite Club, Najafabad Branch, Islamic Azad University, Najafabad, Iran
Hossein Edris Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156–83111, Iran
1
Corresponding Author:
Tel: +98 911 2124023 Email: [email protected]; [email protected]; [email protected]
1
Abstract In the present study, the effect of four-layer accumulative roll bonding (ARB) process at room temperature on the microstructure and mechanical properties of steel-based composite was investigated. Microstructural observations were done by scanning transmission electron microscopy (STEM). Also, textural evolution during the ARB process was evaluated using X-ray diffraction. It was found that occurrence of discontinuously dynamic recrystallization (DDRX) in the interstitial free (IF) steel during cold plastic deformation is possible. The four-layer ARB process at room temperature in presence of SiC microparticles provided the stored energy required for DDRX in the IF steel with high stacking fault energy (SFE). On the other hand, hindrance of initial grain boundaries to dislocation movement resulted in DDRX grains along initial grain boundaries in the IF steel matrix. Average grain size of the final sample was 73 nm but the microstructure was relatively inhomogeneous. The results also indicated that particle stimulated nucleation (PSN) promoted texture randomization in the steel-based composite. Dislocation density of the samples determined from hardness measurement. The dislocation density increased continually until the dislocation density of the steel-based composite after fourth cycle was about 5.3 times (10.73 × 109 cm-2) higher than that of the initial sample (2.02 × 109 cm-2). Also, significant increase in dislocation density 2
occurred after first cycle and in the final cycles the dislocation density value was saturated. After first cycle, a remarkable improve observed in the yield strength value, from 84 MPa to 682 MPa which is almost 8.1 times greater than that of the initial sample. After final cycle, the yield strength value increased to 1061 MPa. Finally, the contribution of individual mechanisms such as the grain refinement, dislocation, second phase, and precipitation in strengthening of the IF steel was evaluated. The contribution of grain refinement and precipitation to the improvement in yield strength was maximum (~67%) and minimum (~3.2%), respectively.
Keywords:
Nanostructured
materials;
Steel-based
composite;
Accumulative roll bonding
1. Introduction Nanostructured and ultra-fine grained (UFG) materials have gained pronounced interest in the scientific community as they show extraordinary physical and mechanical properties [1–3]. The processing of nanostructured and UFG metallic materials using severe plastic deformation (SPD) is a new and promising method of enhancing the properties of metals and alloys for advanced structural and functional applications [4–6]. SPD processing is attractive because the straining is 3
practically unlimited due to the unchanging sample geometry and shape [7]. Up to now, a number of different SPD processing techniques have been developed, such as equal channel angular pressing (ECAP), high pressure torsion (HPT) and accumulative roll bonding (ARB) [1–7]. Among these processes, ECAP and ARB are two of the most frequently used SPD techniques. Accumulative roll bonding was first introduced by Saito et al. [8,9] and is considered to be one of the most promising methods for manufacturing nanostructured and UFG sheet materials. The evolution of microstructures during ARB process were studied for several materials such as composite [10–15], aluminum [8,9,16,17], copper [18], Brass [19], and steel [20–23]. So far, many efforts have been made to achieve nanostructure in the interstitial free (IF) steel but had none of the researchers have managed to create a nanostructure in the IF steel. The minimum grain size of severe plastic deformation methods applied on the IF steel by different researchers was about 300 nm [22– 25]. To achieve nanostructure in the IF steel by ARB process, what should we do? One of the strategies that can be effective is preventing the easy movement of dislocations and occurrence of recovery by second-phase particles. Recently, the present authors have been investigated the mechanical properties of IF steel/SiC composite fabricated by ARB process [26]. By a careful examination of the published literature on the 4
ARB process, it can be noted that no information related to the microstructural parameters that developed in IF steel materials fabricated via the ARB process in the presence of silicon carbide microparticles are available. In order to take full advantage of this emerging technology, this work was undertaken to investigate profoundly the microstructure evolution of IF steel/SiC composite strips processed by ARB process using scanning transmission electron microscopy (STEM) technique.
2. Experimental procedure The materials used in this study were fully annealed strips of interstitial free steel (specifications are given in Table 1) and SiC microparticles (50 µm). The microstructure of the IF steel is shown in Fig. 1. Four strips of 150 mm × 50 mm × 0.7 mm were degreased via acetone and scratch brushed with a stainless steel wire brush 0.25 mm in diameter. After the surface preparation, SiC microparticles were dispersed between the four strips. To achieve a uniform dispersion of SiC particles between IF steel strips, an acetone-based suspension was prepared. After surface preparation, the SiC particles in acetone were sprayed between the four strips with an atomizer. Then, SiC particles were deposited and the acetone evaporated in air, so that the brushed surfaces of strips were uniformly covered with SiC particles. The strips were then stacked over 5
each other and fastened at both ends by steel wires. Attention was also paid to proper alignment of the four strip surfaces prior to rolling. The roll bonding process was carried out with no lubrication and with an amount of thickness reduction equal to 75% corresponding to a von Mises equivalent strain evM of 1.6 per cycle (first step). Then, the roll bonded strips were cut into four pieces. Then, to achieve a uniform distribution of SiC particles in the IF steel matrix, the above procedure was repeated again up to fourth cycle without adding reinforcement particles (second step). The schematic illustration of the ARB process for fabrication of steel composite is shown in Fig. 2. Initial sample for scanning electron microscopy (SEM) observations was cut from the strips and this was mounted in bakelite. Then, the sample was polished using 80–4000 grit water-proof SiC paper. Finally, the polishing was finished on a cloth using alumina paste of 3 µm. Scanning electron microscopy PHILIPS XL30 was used. The microstructural observations were performed using scanning transmission electron microscopy (STEM). Thin foils were prepared with electropolishing conducted at –30 °C using 60 V in a 5% perchloric acid/95% methanol solution. Multiple disc samples with 3 mm diameter were separated from thin foils. Then, the discs were prepared using a low angle Ion Milling System from Fischione Model 1010 with 5 kV operating voltage, 5 mA current, 2.5 h time duration, and angle of 10° conducted at 6
-40 °C. Hitachi S-4800 field emission scanning electron microscope was used. The macrotexture observations were performed on the section which is perpendicular to the transverse direction of rolling (RD–TD plane). Texture was measured on an area of 25 mm × 15 mm of the samples at the depth corresponding to quarter of the thickness. These samples were prepared
by
the
standard
metallographic
procedures.
Texture
measurements were made by recording pole figures of (110) and (200) planes by X–ray diffraction. The experimental pole figure and inverse pole figure data were used to calculate orientation distribution function (ODF) by TexTools software using two measured pole figures. Vickers microhardness was measured according to the ASTM E384-11e1 standard. The surfaces used for indentation testing were ground with SiC papers and polished with a sequence of alumina particles suspensions. Vickers indentation tests were performed using loads in the range 0.01 to 2 N. The tensile test samples were machined from the ARB-processed sheets, according to the ASTM: E8M tensile sample, oriented along the rolling direction. The gauge width and length of the tensile test samples were 6 and 25 mm, respectively. The tensile tests were conducted at room temperature on a Hounsfield H50KS testing machine at an initial strain rate of 1.67 × 10-4 s-1. 7
3. Results and discussion STEM micrographs from the RD–TD plane of steel-based composite produced by ARB process after the first, second, third, and fourth cycles are shown in Fig. 3. As seen in Fig 3(a), the dislocations arranged in different geometries, such as tangling of multiple dislocations, polygonized dislocation walls that form subgrain boundaries, subgrain boundary formed by dislocation rearrangement, and several straight and semicircle dislocations in coarse grains. Relatively high dislocation density can be seen as resulting from one cycle. It should be noted that the SiC microparticles present in the IF steel matrix are responsible for an increase of the dislocation density. These dislocations are generated at
the
reinforcement/matrix
interface
to
accommodate
strain
incompatibility between the two phases. According to Fig. 3(a), beginning stage of subgrain formation can be seen from the changes of contrast. During plastic deformation, a high SFE such as IF steel (200 mJ/m2) strongly confines the separation of partial dislocations that facilitates cross-slip to form three dimensional dislocation substructures like cells or subgrains [27,28]. The average grain size of the sample after the first cycle is 1 µm while that of the initial sample was 25 µm. This indicates that the first cycle of ARB process has a remarkable effect on the grain size of IF steel. However, the distribution of grain size is non8
uniform
because
the
plastic
deformation
of
ARB
process
is
inhomogeneous. It should be noted that there is a very low density of dislocations in the ultrafine grain region as compared with in the coarse grain region. This observation can be rationalized as follows. It is difficult for the ultrafine grains to contain dislocation tangling because when the grain size is below a critical value, the mean free path of dislocations was limited only by its grain boundaries. It should be emphasized that the formation of the ultrafine grains is rather inhomogeneous and depends on the area of observation, in particular in the highly strained materials. The average grain size for two-cycle sample is 700 nm. As shown in Figure 3(b), after the second cycle, some important events occurred in the microstructure of the material. On the one hand, the presence of subgrains that exist within the grains indicates that the continuous dynamic recrystallization (CDRX) mechanism happened. On the other hand, particle stimulated nucleation (PSN) and discontinuous dynamic recrystallization (DDRX) occurred in the vicinity of silicon carbide microparticles.
Dynamic
recrystallization
mechanisms
can
be
distinguished as either continuous or discontinuous [27,28]. The CDRX is partly a dynamic recovery process. It is based on continuous absorption of dislocations in subgrain boundaries, followed by the formation of new grains separated by high angle boundaries. In other 9
words, the development of CDRX is associated with the formation of subgrains with low angle grain boundaries, which are then transformed to grains with high angle boundaries during straining. In contrast, in the DDRX, new grains are formed, usually by conventional nucleation and nucleus growth. For the DDRX process, grain boundary bulging is considered as the initial step for the nucleation of recrystallized grains [27,28]. After the nucleation step, the new DDRX grains grow to use up the deformed matrix. However, for bulging to take place, subgrains near the pre-existing grain boundaries should be well developed in order to provide the required driving force for local migration of the grain boundary [27,28]. In general, dislocations will remain in the interior of recrystallized grains in a CDRX mechanism, whereas DDRX removes dislocations through the sweeping action of high angle grain boundaries. It should be noted that in the hot working of metals with high stacking fault energy (SFE), the occurrence of DDRX has been widely recognized and studied for austenitic steel, copper, nickel, and lead, as well as others [27–34]. However, in the case of aluminum and ferritic steel, reports on the occurrence of DDRX are not as common. The SFE, which is related to the atomic bonding in the material, determines the extent to which unit dislocations
dissociate into
the partial dislocations. Such
dissociation, is promoted by a low value of γSFE, hinders the climb and cross slip of dislocations, which are basic mechanisms responsible for 10
recovery. Stacking fault energy of IF steel is very high (200 mJ/m2). It is well-known that DDRX behavior depends on the SFE and DDRX tends to occur more easily in metals with lower SFE. There is no reports in the literature about the occurrence of DDRX in the IF steel during room temperature deformation. It is believed that rapid dynamic recovery prevents the accumulation of sufficient density of dislocations that is necessary to sustain DDRX in the IF steel [27,28]. In fact, the IF steel recover rapidly during deformation because of its high SFE. Consequently, the stored energy required for DDRX is not achieved. Therefore, dynamic recovery (DRV) and CDRX is usually the only forms of restoration mechanisms during room temperature deformation of the IF steel. But, the present work indicated that the occurrence of DDRX in the IF steel during cold plastic deformation is possible. To the best of the authors’ knowledge, this work is the first to report occurrence of DDRX by ARB at room temperature in IF steel. This phenomenon is due to the presence of SiC microparticles in the IF steel. It is worth noticing that the presence of large (> 1 µm) particles could stimulate the formation of new grains during deformation [27,28]. Previous studies [27,28,35–42] has shown that a region of high dislocation density and large orientation gradient can be formed in the vicinity of microparticles during the deformation. This region is called particle deformation zone (PDZ), which is an ideal site for the 11
development of a recrystallization nucleus through the mechanism of PSN. PSN occurs in the deformation zones around the constituent particles and in three stages [27,28,35–42]: (a) formation of a nucleus from
pre-existing
subgrains
in
the
deformation
zone;
(b)
recrystallization of the deformation zone at the particle; and (c) growth of recrystallized grains beyond the deformation zone. The mechanism of PSN is rapid subboundary migration in the deformation zone that necessarily forms around large (> 1 µm diameter) hard particles during deformation, leading to the creation of new high angle grain boundaries [27,28,35–42]. The deformation zone is characterized by lattice rotation, and hence contains a misorientation gradient. For PSN to occur, the accumulation of misorientation by the rapid subboundary migration must be sufficient to generate the necessary high angle grain boundary. This forms a new grain nucleus, which may then either stagnate, or grow to produce a recrystallized grain. It should be noted that to begin DDRX, the dislocations should accumulate at the particles to form deformation zones, and the accumulated energy should be greater than the critical strain energy required for DDRX. During deformation, dislocations will also be removed by DRV. DDRX in the IF steel will occur when the increase in strain energy due to accumulation of dislocations is greater than the removal of strain energy that occurs through DRV and when the critical strain energy is reached. As a result, four-layer ARB process at 12
room temperature in presence of SiC microparticles can provide the stored energy required for DDRX in the IF steel with high SFE. Fig. 3(c) shows the mean grain size after the third cycle is 350 nm. As can be seen, some DDRX grains formed at pre-existing grain boundaries. It can be concluded that for steel-based composite, grain boundaries can also hinder the movement of dislocations, so dislocations pile up in the vicinity of the initial grain boundaries during ARB process. This can cause initiation of DDRX at grain boundaries through initial grain boundary bulging. A schematic of the initial grain boundary bulging mechanism is shown in Fig. 4. Bulging is known as the classical mechanism for DDRX and occurs when areas of a grain boundary are pinned by low angle boundaries or precipitates [27,28]. It should be noted that the IF steels have a single phase ferritic microstructure with decreased amount of interstitial carbon and nitrogen in the ferrite, which is achieved by microalloying elements of Ti and/or Nb and formation of nanocarbides and nanonitrides. These precipitates are shown in Fig. 5. The nanocarbides and nanonitrides can be effective on the bulging mechanism. As deformation continues, grain boundary shearing leads to inhomogeneous local strain gradients along boundaries, causing initial grain boundaries to bow out (appearing optically as serrated grain boundaries) from local strain induced boundary migrations [27,28]. The new bulged regions of the boundary are dislocation free and become 13
DDRX nuclei. Therefore, the hindrance of the initial grain boundaries to dislocation movement can result in DDRX grains along initial grain boundaries in the IF steel matrix. However, it should be noted that DDRX started near SiC microparticles rather than grain boundaries in this study. Furthermore, the area fractions of DDRX near particles were much larger than those at grain boundaries. After the final cycle (Fig. 3(d)), the average grain size is 73 nm. It is clear that the microstructure is relatively inhomogeneous. Maximum and minimum grain size for four-cycle sample is 250 and 25 nm, respectively. This is due to the presence of silicon carbide microparticles in the IF steel matrix. Since the size of the SiC particles is in the micrometer range and its volume fraction is low (2 Vol.%), therefore, the whole microstructure cannot be influenced by the presence of particles. As mentioned before, in the steel-based composite the hard SiC microparticles rotated relative to IF steel matrix and caused a deformation zone around them. As shown in Fig. 6, in the regions including the particles, the rotation of silicon carbide microparticles made IF steel matrix around particles undergo longer deformation path and undertake larger effective strain compared to the regions without particles. In fact, the effect of SiC particles on the IF steel matrix (formation of PDZ and occurrence of PSN) acted repeatedly during the ARB process and led to a quickly decreasing grain size in the IF steel 14
matrix of the composite. It should be noted that the PDZ can extend to a distance of about a diameter of the particle from the surface of the particle [27]. PDZs can overlap in particle aggregations because the distances between clustered particles are much smaller than a diameter of the particle. However, in this study, because the silicon carbide microparticle spacing in steel-based composite isn’t comparable with the SiC particle size (due to low volume fraction of SiC particles), the deformation zone around particles will not cover the whole IF steel matrix, leading to relatively inhomogeneous deformation microstructure in the IF steel matrix. Previous researchers have reported that the final average grain size of IF steel produced by the ARB process was about 300 nm [1,9,22–25]. In the present work, the average grain size of the IF steel is 73 nm. This difference must be related to the presence of silicon carbide microparticles in the IF steel matrix. It can be concluded that the role of SiC particles in preventing the dislocation motion is very crucial. In fact, in the presence of particles, the stored energy and therefore the driving force for CDRX and DDRX increased, and also the large particles act as nucleation sites through PSN mechanism. One important issue would be to figure out, is the PSN mechanism can affect the texture of the produced materials? Figs. 7 and 8 demonstrate the {200} pole figures and inverse pole figures of ARB-processed steel15
based composite, respectively. After the first cycle, the {011}<100>, {111}<112>, {111}<110>, and {011}<211> components increased. For onecycle sample, the {011}<100> orientation with the maximum intensity of 2.6×R was dominant. After the second cycle, the overall intensity of the components indicated a decrease compared to the previous cycle. For three-cycle sample, all texture components weakened and the {111}<112> with the maximum intensity of 1.5×R was strongest orientation. Figs. 7 and 8 indicated that after second and third cycles, texture weakening occurred. This is due to occurrence of PSN mechanism. Because of the local constraints imposed by the non-deformable SiC microparticles, these regions undergo different strains and will undergo different crystallographic rotations to the material remote from the particles. PSN hence promotes texture randomization in the steel-based composite. Recently, Kobayashi et al. [43] expressed that the presence of nondeformable particles within the 45 deg ND-rotated-cube orientation changes the plane strain deformation field promoting a favorable condition for shear at the vicinity of the particle. This new deformation mode leads to the activation of slip system combinations resulting in the local lattice rotation about the TD. In summary, the very fine grain size (73 nm) is the result of both the high stored energy from four-layer ARB process at room temperature and the
16
presence of SiC microparticles. These particles act as sites for PSN, which leads to the relative random texture. Dislocation density is a critical variable that determines dislocation mobility, strength and ductility. Dislocation density is closely related to the ability of a given material to endure plastic deformation. Recently, an alternative way to measure dislocation density developed by Graca et al. [44] based on the application of an ISE model [45]. They mentioned that the dislocation density of the materials could be estimated from indentation hardness measurement. The hardness values increase as the indentation depth decreases. The relation between the hardness value and indentation depth is indicated as following equation [45]: 2
H * 1 =1+ h h H0
(1)
where H0 is the hardness in the limitation of infinite depth (bulk hardness), h* is a characteristic length and H is the hardness value corresponding to indentation depth h. The correlation between the dislocation density stored in the lattice and the characteristic length is as following [44]:
ρ=
3 1 tan 2 θ 2 f 3 bh *
(2)
where b is the Burgers vector of the dislocations, ρ is the dislocation density, θ is the angle between the surfaces of the material and the 17
indenter, and f is a correction factor for the size of the plastic zone. It should be noted that the indentation depth, h, was calculated by [46,47]: h=
d 2 2 tan
θ 2
(3)
where d is the indentation diameter. In the present work θ=22°, f=1.9, and b=0.248 nm [48,49]. Introducing these values in Eq. (2) a dislocation density is obtained. Fig. 9 demonstrates the value of h* in the ARB-processed samples through 0, 1, 2, 3, and 4 cycles which are measured to be 7125.99, 2083.14, 1597.61, 1385.42, and 1341.52 nm, respectively determined by fitting Eq. (1) to experimental data. The dislocation density of samples versus number of cycles calculated from Eq. (2) is shown in Fig. 10. It is obvious that with increasing the number of ARB cycles, the dislocation density increases. The dislocation density increased from 2.02 × 109 cm-2 (for the initial sample) to 6.91 × 109 cm-2 for one-cycle sample, registering 242% increase. It can be concluded that the first cycle of ARB process has remarkable effect on the dislocation density. The values of dislocation density are 9.01 × 109 and 10.39 × 109 cm-2 after second and third cycles, respectively. The dislocation density increased continually until the dislocation density of the steel-based composite after fourth cycle was about 5.3 times (10.73 × 109 cm-2) higher than that of the initial sample. Regarding Fig. 10, the value of dislocation density first rapidly increased, then dwindled, and finally 18
saturated by further ARB cycles. The saturation of dislocation density occurs because the material reached to the steady-state density of dislocation. The steady-state density of dislocations is attributed to dynamic balance between dislocation generation during ARB process and annihilation by occurrence of dynamic restoration phenomena. Therefore, it can be concluded that occurrence of DRV, CDRX, and DDRX concur to keep the dislocation density constant for high strains. Engineering stress–strain curves of the initial sample and the ARBprocessed steel-based composite samples are shown in Fig. 11. Also, the variations of yield strength for steel-based composite during sequential ARB cycles is presented in Fig. 12. After first cycle, a remarkable improve can be observed in the yield strength value, from 84 MPa to 682 MPa which is almost 8.1 times greater than that of the initial sample. The significant increase in yield strength at the first cycle can be attributed to strain hardening (i.e., increase in density of dislocations and their interactions) and formation of subgrains. The yield strength is 794 and 969 MPa after second and third ARB cycles, respectively. After final cycle, the yield strength value of steel-based composite increased to 1061 MPa. It can be said that that yield strength variations in the ARBprocessed
materials
are
governed
by
two
main
strengthening
mechanisms: strain hardening and grain refinement. In the early stages of ARB process, strain hardening plays a main role in increasing the 19
yield strength, while the formation of submicron subgrains or dislocation cells also contributes to strength. But at higher cycles, higher yield strength is achieved by grain refinement as the ARB cycles increase. This result is consistent with the discussion of Figs. 3 and 10. It should be noted that the SiC microparticles also have an outstanding effect on the tensile behavior of steel-based composite. There are two different strengthening mechanisms that are typically associated with metal matrix composites: direct strengthening resulting from load transfer from the metal matrix to the reinforcing particle and indirect strengthening resulting from the influence of reinforcement on matrix microstructure or deformation mode, such as dislocation strengthening induced by the deformation mismatch between the reinforcement and the matrix. The steel/SiC interface plays an important role in determining the mechanical response of the steel-based composite, and its influence can be examined in the context of a load transfer mechanism.
Under
these
conditions,
the
contribution
of
the
microstructure to the mechanical properties can be estimated on the basis of a simple rule-of-mixtures [50–52]:
σ = σ m (1 − f ) + σ r f
(4)
where σm is the yield strength of the metal matrix, σr is the yield strength of the reinforcement phase, and f is the volume fraction of the reinforcement phase. In the present work σ=1061 MPa, f=0.02, and 20
σr=3440 MPa [53,54]. Therefore, the yield strength of the IF steel matrix (σm) is calculated as 1012.5 MPa. It should be noted that the contribution of reinforcement particles to the improvement in yield strength of the IF steel through its direct strengthening effect is equal to 68.8 MPa. On the other hand the IF steel matrix can be strengthened by three different
mechanisms
including
grain
refinement
strengthening,
dislocation strengthening, and precipitation strengthening. It is clear that σm must be a direct consequence of the aforementioned strengthening effects. With this in mind, an attempt has been made to separate the influence of certain parameters on strengthening of composites. The contribution of various strengthening mechanisms to the improvement in σm is discussed as follows. σm can be estimated by the following equation: (5)
σ m = σ 0 + ∆σ g + ∆σ dis + ∆σ p
where σ0 is the yield strength of undeformed pure single-crystal IF steel (σ0=30 MPa [48]), and ∆σg, ∆σdis, and ∆σp are the increases in the yield strength caused by grain refinement, dislocation, and precipitation, respectively. The contribution of the grain refinement ∆σg is first evaluated. ∆σg can be calculated by the following equation [48,49]: ∆σ g = kd
−
1 2
(6)
21
where k is the Hall–Petch slope (k=200 MPa µm0.5 for IF steel [49]) and d is the average grain size (d=73 nm). The contribution of grain refinement to the improvement in yield strength of the IF steel matrix is then calculated as 707.1 MPa. Next, the contribution of dislocation strengthening is analyzed. ∆σdis can be calculated by the following equation [48,49]: ∆σ dis = M αGb ρ
1 2
(7)
where M=3.2 [49] is the Taylor factor, α=0.33 [48] is a constant, G=77 GPa [49] is shear modulus, b=0.248 nm [48] is Burgers vector, and ρ is the dislocation density (10.73 × 109 cm-2). The contribution of dislocation strengthening to the improvement in yield strength of the IF steel matrix is then calculated as 208.9 MPa. Finally, the contribution of precipitation strengthening of nanocarbides and nanonitrides is determined. Based on 1012.5 MPa of the total yield strength of the IF steel matrix and according to Eq. (5), the contribution of precipitation to the improvement is then calculated as 33.4 MPa. Fig. 13 shows the contribution of various strengthening mechanisms to the improvement in yield strength of ARB-processed steel-based composite. As seen, the contributions of different strengthening methods are different in the case of the steel-based composite. The contribution of grain refinement and precipitation to the improvement in yield strength is maximum (~67%)
22
and minimum (~3.2%), respectively. It can be concluded that the predominant method for the improvement of IF steel strength is the grain refining technology.
4. Conclusions In this research, the effect of four-layer accumulative roll bonding (ARB) process at room temperature on the microstructure and mechanical properties of steel-based composite was investigated. The conclusions drawn from the results can be summarized as follows: 1) Occurrence of DDRX in the IF steel during cold plastic deformation was possible. This phenomenon was due to the presence of SiC microparticles in the IF steel. The four-layer ARB process at room temperature in presence of SiC microparticles provided the stored energy required for DDRX in the IF steel with high SFE. 2) Average grain size of the final sample was 73 nm but the microstructure was relatively inhomogeneous. 3) In the presence of particles, the stored energy and therefore the driving force for CDRX and DDRX increased, and also the large particles acted as nucleation sites through PSN mechanism. 4) After second and third cycles, texture weakening occurred. This is due to occurrence of PSN mechanism. Because of the local 23
constraints imposed by the non-deformable SiC microparticles, these regions undergo different strains and will undergo different crystallographic rotations to the material remote from the particles. 5) First cycle of ARB process had remarkable effect on the dislocation density. The dislocation density increased continually until the dislocation density of the steel-based composite after fourth cycle was about 5.3 times (10.73 × 109 cm-2) higher than that of the initial sample (2.02 × 109 cm-2). The value of dislocation density first rapidly increased, then dwindled, and finally saturated by further ARB cycles. It can be concluded that occurrence of DRV, CDRX, and DDRX concur to keep the dislocation density constant for high strains. 6) After first cycle, a remarkable improve observed in the yield strength value, from 84 MPa to 682 MPa which is almost 8.1 times greater than that of the initial sample. After final cycle, the yield strength value of steel-based composite increased to 1061 MPa. 7) The contribution of grain refinement and precipitation to the improvement in yield strength was maximum (~67%) and minimum (~3.2%), respectively. The predominant method for the improvement of IF steel strength was the grain refining technology. 24
Acknowledgement The corresponding author gratefully acknowledges Mrs. Mahjoobeh Hatef.
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33
Fig. 1. Microstructure of the IF steel at two magnifications.
34
Fig. 2. Schematic illustration of ARB process for fabrication of steel composite.
35
36
37
38
Fig. 3. STEM micrographs from RD–TD plane of samples after (a) first, (b) second, (c) third, and (d) fourth cycles.
39
Fig. 4. Schematic illustration of: (a) boundaries dragged by precipitates and bulging driven by locally concentrated dislocations, and (b) boundaries dragged and bulging driven by low angle boundaries.
40
Fig. 5. The nanoprecipitates in the IF steel microstructure.
41
Fig. 6. Schematic illustration of the deformation path for: (a) regions without particles showing relative straight sliding path, and (b) regions including particles showing longer deformation path due to the rotation of silicon carbide microparticles. 42
43
Fig. 7. {200} pole figures (a) before ARB and after (b) 1, (c) 2, and (d) 3 ARB cycles. 44
45
Fig. 8. Inverse pole figures (a) before ARB and after (b) 1, (c) 2, and (d) 3 ARB cycles.
46
Fig. 9. The h* values for ARB-processed samples calculated by Eq. (1).
47
Fig. 10. The dislocation densities for ARB-processed samples calculated by Eq. (2) based on h* values of Fig. 9.
48
Fig. 11. Engineering stress–strain curves of initial and ARB-processed samples.
49
Fig. 12. Variation of yield strength in ARB-processed samples for different cycles.
50
Fig. 13. Contribution of strengthening mechanisms to the improvement in yield strength of ARB-processed steel-based composite.
51
Tables:
Table 1. Chemical composition of IF steel (in wt.%). C
N
Si
Mn
Cu
Ni
Ti
Fe
0.002
0.004
0.01
0.14
0.01
0.018
0.055
Bal.
52
PII: DOI: Reference:
S0921-5093(15)00547-X http://dx.doi.org/10.1016/j.msea.2015.05.025 MSA32353
To appear in: Materials Science & Engineering A Received date: 25 February 2015 Revised date: 7 May 2015 Accepted date: 8 May 2015 Cite this article as: Roohollah Jamaati, Mohammad Reza Toroghinejad, Sajjad Amirkhanlou and Hossein Edris, Microstructural evolution of nanostructured steel-based composite fabricated by accumulative roll bonding, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.05.025 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructural evolution of nanostructured steelbased composite fabricated by accumulative roll bonding
Roohollah Jamaati 1 Department of Mechanical Engineering, Babol Noshirvani University of Technology, Babol, Iran
Mohammad Reza Toroghinejad Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156–83111, Iran
Sajjad Amirkhanlou Young Researchers and Elite Club, Najafabad Branch, Islamic Azad University, Najafabad, Iran
Hossein Edris Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156–83111, Iran
1
Corresponding Author:
Tel: +98 911 2124023 Email: [email protected]; [email protected]; [email protected]
1
Abstract In the present study, the effect of four-layer accumulative roll bonding (ARB) process at room temperature on the microstructure and mechanical properties of steel-based composite was investigated. Microstructural observations were done by scanning transmission electron microscopy (STEM). Also, textural evolution during the ARB process was evaluated using X-ray diffraction. It was found that occurrence of discontinuously dynamic recrystallization (DDRX) in the interstitial free (IF) steel during cold plastic deformation is possible. The four-layer ARB process at room temperature in presence of SiC microparticles provided the stored energy required for DDRX in the IF steel with high stacking fault energy (SFE). On the other hand, hindrance of initial grain boundaries to dislocation movement resulted in DDRX grains along initial grain boundaries in the IF steel matrix. Average grain size of the final sample was 73 nm but the microstructure was relatively inhomogeneous. The results also indicated that particle stimulated nucleation (PSN) promoted texture randomization in the steel-based composite. Dislocation density of the samples determined from hardness measurement. The dislocation density increased continually until the dislocation density of the steel-based composite after fourth cycle was about 5.3 times (10.73 × 109 cm-2) higher than that of the initial sample (2.02 × 109 cm-2). Also, significant increase in dislocation density 2
occurred after first cycle and in the final cycles the dislocation density value was saturated. After first cycle, a remarkable improve observed in the yield strength value, from 84 MPa to 682 MPa which is almost 8.1 times greater than that of the initial sample. After final cycle, the yield strength value increased to 1061 MPa. Finally, the contribution of individual mechanisms such as the grain refinement, dislocation, second phase, and precipitation in strengthening of the IF steel was evaluated. The contribution of grain refinement and precipitation to the improvement in yield strength was maximum (~67%) and minimum (~3.2%), respectively.
Keywords:
Nanostructured
materials;
Steel-based
composite;
Accumulative roll bonding
1. Introduction Nanostructured and ultra-fine grained (UFG) materials have gained pronounced interest in the scientific community as they show extraordinary physical and mechanical properties [1–3]. The processing of nanostructured and UFG metallic materials using severe plastic deformation (SPD) is a new and promising method of enhancing the properties of metals and alloys for advanced structural and functional applications [4–6]. SPD processing is attractive because the straining is 3
practically unlimited due to the unchanging sample geometry and shape [7]. Up to now, a number of different SPD processing techniques have been developed, such as equal channel angular pressing (ECAP), high pressure torsion (HPT) and accumulative roll bonding (ARB) [1–7]. Among these processes, ECAP and ARB are two of the most frequently used SPD techniques. Accumulative roll bonding was first introduced by Saito et al. [8,9] and is considered to be one of the most promising methods for manufacturing nanostructured and UFG sheet materials. The evolution of microstructures during ARB process were studied for several materials such as composite [10–15], aluminum [8,9,16,17], copper [18], Brass [19], and steel [20–23]. So far, many efforts have been made to achieve nanostructure in the interstitial free (IF) steel but had none of the researchers have managed to create a nanostructure in the IF steel. The minimum grain size of severe plastic deformation methods applied on the IF steel by different researchers was about 300 nm [22– 25]. To achieve nanostructure in the IF steel by ARB process, what should we do? One of the strategies that can be effective is preventing the easy movement of dislocations and occurrence of recovery by second-phase particles. Recently, the present authors have been investigated the mechanical properties of IF steel/SiC composite fabricated by ARB process [26]. By a careful examination of the published literature on the 4
ARB process, it can be noted that no information related to the microstructural parameters that developed in IF steel materials fabricated via the ARB process in the presence of silicon carbide microparticles are available. In order to take full advantage of this emerging technology, this work was undertaken to investigate profoundly the microstructure evolution of IF steel/SiC composite strips processed by ARB process using scanning transmission electron microscopy (STEM) technique.
2. Experimental procedure The materials used in this study were fully annealed strips of interstitial free steel (specifications are given in Table 1) and SiC microparticles (50 µm). The microstructure of the IF steel is shown in Fig. 1. Four strips of 150 mm × 50 mm × 0.7 mm were degreased via acetone and scratch brushed with a stainless steel wire brush 0.25 mm in diameter. After the surface preparation, SiC microparticles were dispersed between the four strips. To achieve a uniform dispersion of SiC particles between IF steel strips, an acetone-based suspension was prepared. After surface preparation, the SiC particles in acetone were sprayed between the four strips with an atomizer. Then, SiC particles were deposited and the acetone evaporated in air, so that the brushed surfaces of strips were uniformly covered with SiC particles. The strips were then stacked over 5
each other and fastened at both ends by steel wires. Attention was also paid to proper alignment of the four strip surfaces prior to rolling. The roll bonding process was carried out with no lubrication and with an amount of thickness reduction equal to 75% corresponding to a von Mises equivalent strain evM of 1.6 per cycle (first step). Then, the roll bonded strips were cut into four pieces. Then, to achieve a uniform distribution of SiC particles in the IF steel matrix, the above procedure was repeated again up to fourth cycle without adding reinforcement particles (second step). The schematic illustration of the ARB process for fabrication of steel composite is shown in Fig. 2. Initial sample for scanning electron microscopy (SEM) observations was cut from the strips and this was mounted in bakelite. Then, the sample was polished using 80–4000 grit water-proof SiC paper. Finally, the polishing was finished on a cloth using alumina paste of 3 µm. Scanning electron microscopy PHILIPS XL30 was used. The microstructural observations were performed using scanning transmission electron microscopy (STEM). Thin foils were prepared with electropolishing conducted at –30 °C using 60 V in a 5% perchloric acid/95% methanol solution. Multiple disc samples with 3 mm diameter were separated from thin foils. Then, the discs were prepared using a low angle Ion Milling System from Fischione Model 1010 with 5 kV operating voltage, 5 mA current, 2.5 h time duration, and angle of 10° conducted at 6
-40 °C. Hitachi S-4800 field emission scanning electron microscope was used. The macrotexture observations were performed on the section which is perpendicular to the transverse direction of rolling (RD–TD plane). Texture was measured on an area of 25 mm × 15 mm of the samples at the depth corresponding to quarter of the thickness. These samples were prepared
by
the
standard
metallographic
procedures.
Texture
measurements were made by recording pole figures of (110) and (200) planes by X–ray diffraction. The experimental pole figure and inverse pole figure data were used to calculate orientation distribution function (ODF) by TexTools software using two measured pole figures. Vickers microhardness was measured according to the ASTM E384-11e1 standard. The surfaces used for indentation testing were ground with SiC papers and polished with a sequence of alumina particles suspensions. Vickers indentation tests were performed using loads in the range 0.01 to 2 N. The tensile test samples were machined from the ARB-processed sheets, according to the ASTM: E8M tensile sample, oriented along the rolling direction. The gauge width and length of the tensile test samples were 6 and 25 mm, respectively. The tensile tests were conducted at room temperature on a Hounsfield H50KS testing machine at an initial strain rate of 1.67 × 10-4 s-1. 7
3. Results and discussion STEM micrographs from the RD–TD plane of steel-based composite produced by ARB process after the first, second, third, and fourth cycles are shown in Fig. 3. As seen in Fig 3(a), the dislocations arranged in different geometries, such as tangling of multiple dislocations, polygonized dislocation walls that form subgrain boundaries, subgrain boundary formed by dislocation rearrangement, and several straight and semicircle dislocations in coarse grains. Relatively high dislocation density can be seen as resulting from one cycle. It should be noted that the SiC microparticles present in the IF steel matrix are responsible for an increase of the dislocation density. These dislocations are generated at
the
reinforcement/matrix
interface
to
accommodate
strain
incompatibility between the two phases. According to Fig. 3(a), beginning stage of subgrain formation can be seen from the changes of contrast. During plastic deformation, a high SFE such as IF steel (200 mJ/m2) strongly confines the separation of partial dislocations that facilitates cross-slip to form three dimensional dislocation substructures like cells or subgrains [27,28]. The average grain size of the sample after the first cycle is 1 µm while that of the initial sample was 25 µm. This indicates that the first cycle of ARB process has a remarkable effect on the grain size of IF steel. However, the distribution of grain size is non8
uniform
because
the
plastic
deformation
of
ARB
process
is
inhomogeneous. It should be noted that there is a very low density of dislocations in the ultrafine grain region as compared with in the coarse grain region. This observation can be rationalized as follows. It is difficult for the ultrafine grains to contain dislocation tangling because when the grain size is below a critical value, the mean free path of dislocations was limited only by its grain boundaries. It should be emphasized that the formation of the ultrafine grains is rather inhomogeneous and depends on the area of observation, in particular in the highly strained materials. The average grain size for two-cycle sample is 700 nm. As shown in Figure 3(b), after the second cycle, some important events occurred in the microstructure of the material. On the one hand, the presence of subgrains that exist within the grains indicates that the continuous dynamic recrystallization (CDRX) mechanism happened. On the other hand, particle stimulated nucleation (PSN) and discontinuous dynamic recrystallization (DDRX) occurred in the vicinity of silicon carbide microparticles.
Dynamic
recrystallization
mechanisms
can
be
distinguished as either continuous or discontinuous [27,28]. The CDRX is partly a dynamic recovery process. It is based on continuous absorption of dislocations in subgrain boundaries, followed by the formation of new grains separated by high angle boundaries. In other 9
words, the development of CDRX is associated with the formation of subgrains with low angle grain boundaries, which are then transformed to grains with high angle boundaries during straining. In contrast, in the DDRX, new grains are formed, usually by conventional nucleation and nucleus growth. For the DDRX process, grain boundary bulging is considered as the initial step for the nucleation of recrystallized grains [27,28]. After the nucleation step, the new DDRX grains grow to use up the deformed matrix. However, for bulging to take place, subgrains near the pre-existing grain boundaries should be well developed in order to provide the required driving force for local migration of the grain boundary [27,28]. In general, dislocations will remain in the interior of recrystallized grains in a CDRX mechanism, whereas DDRX removes dislocations through the sweeping action of high angle grain boundaries. It should be noted that in the hot working of metals with high stacking fault energy (SFE), the occurrence of DDRX has been widely recognized and studied for austenitic steel, copper, nickel, and lead, as well as others [27–34]. However, in the case of aluminum and ferritic steel, reports on the occurrence of DDRX are not as common. The SFE, which is related to the atomic bonding in the material, determines the extent to which unit dislocations
dissociate into
the partial dislocations. Such
dissociation, is promoted by a low value of γSFE, hinders the climb and cross slip of dislocations, which are basic mechanisms responsible for 10
recovery. Stacking fault energy of IF steel is very high (200 mJ/m2). It is well-known that DDRX behavior depends on the SFE and DDRX tends to occur more easily in metals with lower SFE. There is no reports in the literature about the occurrence of DDRX in the IF steel during room temperature deformation. It is believed that rapid dynamic recovery prevents the accumulation of sufficient density of dislocations that is necessary to sustain DDRX in the IF steel [27,28]. In fact, the IF steel recover rapidly during deformation because of its high SFE. Consequently, the stored energy required for DDRX is not achieved. Therefore, dynamic recovery (DRV) and CDRX is usually the only forms of restoration mechanisms during room temperature deformation of the IF steel. But, the present work indicated that the occurrence of DDRX in the IF steel during cold plastic deformation is possible. To the best of the authors’ knowledge, this work is the first to report occurrence of DDRX by ARB at room temperature in IF steel. This phenomenon is due to the presence of SiC microparticles in the IF steel. It is worth noticing that the presence of large (> 1 µm) particles could stimulate the formation of new grains during deformation [27,28]. Previous studies [27,28,35–42] has shown that a region of high dislocation density and large orientation gradient can be formed in the vicinity of microparticles during the deformation. This region is called particle deformation zone (PDZ), which is an ideal site for the 11
development of a recrystallization nucleus through the mechanism of PSN. PSN occurs in the deformation zones around the constituent particles and in three stages [27,28,35–42]: (a) formation of a nucleus from
pre-existing
subgrains
in
the
deformation
zone;
(b)
recrystallization of the deformation zone at the particle; and (c) growth of recrystallized grains beyond the deformation zone. The mechanism of PSN is rapid subboundary migration in the deformation zone that necessarily forms around large (> 1 µm diameter) hard particles during deformation, leading to the creation of new high angle grain boundaries [27,28,35–42]. The deformation zone is characterized by lattice rotation, and hence contains a misorientation gradient. For PSN to occur, the accumulation of misorientation by the rapid subboundary migration must be sufficient to generate the necessary high angle grain boundary. This forms a new grain nucleus, which may then either stagnate, or grow to produce a recrystallized grain. It should be noted that to begin DDRX, the dislocations should accumulate at the particles to form deformation zones, and the accumulated energy should be greater than the critical strain energy required for DDRX. During deformation, dislocations will also be removed by DRV. DDRX in the IF steel will occur when the increase in strain energy due to accumulation of dislocations is greater than the removal of strain energy that occurs through DRV and when the critical strain energy is reached. As a result, four-layer ARB process at 12
room temperature in presence of SiC microparticles can provide the stored energy required for DDRX in the IF steel with high SFE. Fig. 3(c) shows the mean grain size after the third cycle is 350 nm. As can be seen, some DDRX grains formed at pre-existing grain boundaries. It can be concluded that for steel-based composite, grain boundaries can also hinder the movement of dislocations, so dislocations pile up in the vicinity of the initial grain boundaries during ARB process. This can cause initiation of DDRX at grain boundaries through initial grain boundary bulging. A schematic of the initial grain boundary bulging mechanism is shown in Fig. 4. Bulging is known as the classical mechanism for DDRX and occurs when areas of a grain boundary are pinned by low angle boundaries or precipitates [27,28]. It should be noted that the IF steels have a single phase ferritic microstructure with decreased amount of interstitial carbon and nitrogen in the ferrite, which is achieved by microalloying elements of Ti and/or Nb and formation of nanocarbides and nanonitrides. These precipitates are shown in Fig. 5. The nanocarbides and nanonitrides can be effective on the bulging mechanism. As deformation continues, grain boundary shearing leads to inhomogeneous local strain gradients along boundaries, causing initial grain boundaries to bow out (appearing optically as serrated grain boundaries) from local strain induced boundary migrations [27,28]. The new bulged regions of the boundary are dislocation free and become 13
DDRX nuclei. Therefore, the hindrance of the initial grain boundaries to dislocation movement can result in DDRX grains along initial grain boundaries in the IF steel matrix. However, it should be noted that DDRX started near SiC microparticles rather than grain boundaries in this study. Furthermore, the area fractions of DDRX near particles were much larger than those at grain boundaries. After the final cycle (Fig. 3(d)), the average grain size is 73 nm. It is clear that the microstructure is relatively inhomogeneous. Maximum and minimum grain size for four-cycle sample is 250 and 25 nm, respectively. This is due to the presence of silicon carbide microparticles in the IF steel matrix. Since the size of the SiC particles is in the micrometer range and its volume fraction is low (2 Vol.%), therefore, the whole microstructure cannot be influenced by the presence of particles. As mentioned before, in the steel-based composite the hard SiC microparticles rotated relative to IF steel matrix and caused a deformation zone around them. As shown in Fig. 6, in the regions including the particles, the rotation of silicon carbide microparticles made IF steel matrix around particles undergo longer deformation path and undertake larger effective strain compared to the regions without particles. In fact, the effect of SiC particles on the IF steel matrix (formation of PDZ and occurrence of PSN) acted repeatedly during the ARB process and led to a quickly decreasing grain size in the IF steel 14
matrix of the composite. It should be noted that the PDZ can extend to a distance of about a diameter of the particle from the surface of the particle [27]. PDZs can overlap in particle aggregations because the distances between clustered particles are much smaller than a diameter of the particle. However, in this study, because the silicon carbide microparticle spacing in steel-based composite isn’t comparable with the SiC particle size (due to low volume fraction of SiC particles), the deformation zone around particles will not cover the whole IF steel matrix, leading to relatively inhomogeneous deformation microstructure in the IF steel matrix. Previous researchers have reported that the final average grain size of IF steel produced by the ARB process was about 300 nm [1,9,22–25]. In the present work, the average grain size of the IF steel is 73 nm. This difference must be related to the presence of silicon carbide microparticles in the IF steel matrix. It can be concluded that the role of SiC particles in preventing the dislocation motion is very crucial. In fact, in the presence of particles, the stored energy and therefore the driving force for CDRX and DDRX increased, and also the large particles act as nucleation sites through PSN mechanism. One important issue would be to figure out, is the PSN mechanism can affect the texture of the produced materials? Figs. 7 and 8 demonstrate the {200} pole figures and inverse pole figures of ARB-processed steel15
based composite, respectively. After the first cycle, the {011}<100>, {111}<112>, {111}<110>, and {011}<211> components increased. For onecycle sample, the {011}<100> orientation with the maximum intensity of 2.6×R was dominant. After the second cycle, the overall intensity of the components indicated a decrease compared to the previous cycle. For three-cycle sample, all texture components weakened and the {111}<112> with the maximum intensity of 1.5×R was strongest orientation. Figs. 7 and 8 indicated that after second and third cycles, texture weakening occurred. This is due to occurrence of PSN mechanism. Because of the local constraints imposed by the non-deformable SiC microparticles, these regions undergo different strains and will undergo different crystallographic rotations to the material remote from the particles. PSN hence promotes texture randomization in the steel-based composite. Recently, Kobayashi et al. [43] expressed that the presence of nondeformable particles within the 45 deg ND-rotated-cube orientation changes the plane strain deformation field promoting a favorable condition for shear at the vicinity of the particle. This new deformation mode leads to the activation of slip system combinations resulting in the local lattice rotation about the TD. In summary, the very fine grain size (73 nm) is the result of both the high stored energy from four-layer ARB process at room temperature and the
16
presence of SiC microparticles. These particles act as sites for PSN, which leads to the relative random texture. Dislocation density is a critical variable that determines dislocation mobility, strength and ductility. Dislocation density is closely related to the ability of a given material to endure plastic deformation. Recently, an alternative way to measure dislocation density developed by Graca et al. [44] based on the application of an ISE model [45]. They mentioned that the dislocation density of the materials could be estimated from indentation hardness measurement. The hardness values increase as the indentation depth decreases. The relation between the hardness value and indentation depth is indicated as following equation [45]: 2
H * 1 =1+ h h H0
(1)
where H0 is the hardness in the limitation of infinite depth (bulk hardness), h* is a characteristic length and H is the hardness value corresponding to indentation depth h. The correlation between the dislocation density stored in the lattice and the characteristic length is as following [44]:
ρ=
3 1 tan 2 θ 2 f 3 bh *
(2)
where b is the Burgers vector of the dislocations, ρ is the dislocation density, θ is the angle between the surfaces of the material and the 17
indenter, and f is a correction factor for the size of the plastic zone. It should be noted that the indentation depth, h, was calculated by [46,47]: h=
d 2 2 tan
θ 2
(3)
where d is the indentation diameter. In the present work θ=22°, f=1.9, and b=0.248 nm [48,49]. Introducing these values in Eq. (2) a dislocation density is obtained. Fig. 9 demonstrates the value of h* in the ARB-processed samples through 0, 1, 2, 3, and 4 cycles which are measured to be 7125.99, 2083.14, 1597.61, 1385.42, and 1341.52 nm, respectively determined by fitting Eq. (1) to experimental data. The dislocation density of samples versus number of cycles calculated from Eq. (2) is shown in Fig. 10. It is obvious that with increasing the number of ARB cycles, the dislocation density increases. The dislocation density increased from 2.02 × 109 cm-2 (for the initial sample) to 6.91 × 109 cm-2 for one-cycle sample, registering 242% increase. It can be concluded that the first cycle of ARB process has remarkable effect on the dislocation density. The values of dislocation density are 9.01 × 109 and 10.39 × 109 cm-2 after second and third cycles, respectively. The dislocation density increased continually until the dislocation density of the steel-based composite after fourth cycle was about 5.3 times (10.73 × 109 cm-2) higher than that of the initial sample. Regarding Fig. 10, the value of dislocation density first rapidly increased, then dwindled, and finally 18
saturated by further ARB cycles. The saturation of dislocation density occurs because the material reached to the steady-state density of dislocation. The steady-state density of dislocations is attributed to dynamic balance between dislocation generation during ARB process and annihilation by occurrence of dynamic restoration phenomena. Therefore, it can be concluded that occurrence of DRV, CDRX, and DDRX concur to keep the dislocation density constant for high strains. Engineering stress–strain curves of the initial sample and the ARBprocessed steel-based composite samples are shown in Fig. 11. Also, the variations of yield strength for steel-based composite during sequential ARB cycles is presented in Fig. 12. After first cycle, a remarkable improve can be observed in the yield strength value, from 84 MPa to 682 MPa which is almost 8.1 times greater than that of the initial sample. The significant increase in yield strength at the first cycle can be attributed to strain hardening (i.e., increase in density of dislocations and their interactions) and formation of subgrains. The yield strength is 794 and 969 MPa after second and third ARB cycles, respectively. After final cycle, the yield strength value of steel-based composite increased to 1061 MPa. It can be said that that yield strength variations in the ARBprocessed
materials
are
governed
by
two
main
strengthening
mechanisms: strain hardening and grain refinement. In the early stages of ARB process, strain hardening plays a main role in increasing the 19
yield strength, while the formation of submicron subgrains or dislocation cells also contributes to strength. But at higher cycles, higher yield strength is achieved by grain refinement as the ARB cycles increase. This result is consistent with the discussion of Figs. 3 and 10. It should be noted that the SiC microparticles also have an outstanding effect on the tensile behavior of steel-based composite. There are two different strengthening mechanisms that are typically associated with metal matrix composites: direct strengthening resulting from load transfer from the metal matrix to the reinforcing particle and indirect strengthening resulting from the influence of reinforcement on matrix microstructure or deformation mode, such as dislocation strengthening induced by the deformation mismatch between the reinforcement and the matrix. The steel/SiC interface plays an important role in determining the mechanical response of the steel-based composite, and its influence can be examined in the context of a load transfer mechanism.
Under
these
conditions,
the
contribution
of
the
microstructure to the mechanical properties can be estimated on the basis of a simple rule-of-mixtures [50–52]:
σ = σ m (1 − f ) + σ r f
(4)
where σm is the yield strength of the metal matrix, σr is the yield strength of the reinforcement phase, and f is the volume fraction of the reinforcement phase. In the present work σ=1061 MPa, f=0.02, and 20
σr=3440 MPa [53,54]. Therefore, the yield strength of the IF steel matrix (σm) is calculated as 1012.5 MPa. It should be noted that the contribution of reinforcement particles to the improvement in yield strength of the IF steel through its direct strengthening effect is equal to 68.8 MPa. On the other hand the IF steel matrix can be strengthened by three different
mechanisms
including
grain
refinement
strengthening,
dislocation strengthening, and precipitation strengthening. It is clear that σm must be a direct consequence of the aforementioned strengthening effects. With this in mind, an attempt has been made to separate the influence of certain parameters on strengthening of composites. The contribution of various strengthening mechanisms to the improvement in σm is discussed as follows. σm can be estimated by the following equation: (5)
σ m = σ 0 + ∆σ g + ∆σ dis + ∆σ p
where σ0 is the yield strength of undeformed pure single-crystal IF steel (σ0=30 MPa [48]), and ∆σg, ∆σdis, and ∆σp are the increases in the yield strength caused by grain refinement, dislocation, and precipitation, respectively. The contribution of the grain refinement ∆σg is first evaluated. ∆σg can be calculated by the following equation [48,49]: ∆σ g = kd
−
1 2
(6)
21
where k is the Hall–Petch slope (k=200 MPa µm0.5 for IF steel [49]) and d is the average grain size (d=73 nm). The contribution of grain refinement to the improvement in yield strength of the IF steel matrix is then calculated as 707.1 MPa. Next, the contribution of dislocation strengthening is analyzed. ∆σdis can be calculated by the following equation [48,49]: ∆σ dis = M αGb ρ
1 2
(7)
where M=3.2 [49] is the Taylor factor, α=0.33 [48] is a constant, G=77 GPa [49] is shear modulus, b=0.248 nm [48] is Burgers vector, and ρ is the dislocation density (10.73 × 109 cm-2). The contribution of dislocation strengthening to the improvement in yield strength of the IF steel matrix is then calculated as 208.9 MPa. Finally, the contribution of precipitation strengthening of nanocarbides and nanonitrides is determined. Based on 1012.5 MPa of the total yield strength of the IF steel matrix and according to Eq. (5), the contribution of precipitation to the improvement is then calculated as 33.4 MPa. Fig. 13 shows the contribution of various strengthening mechanisms to the improvement in yield strength of ARB-processed steel-based composite. As seen, the contributions of different strengthening methods are different in the case of the steel-based composite. The contribution of grain refinement and precipitation to the improvement in yield strength is maximum (~67%)
22
and minimum (~3.2%), respectively. It can be concluded that the predominant method for the improvement of IF steel strength is the grain refining technology.
4. Conclusions In this research, the effect of four-layer accumulative roll bonding (ARB) process at room temperature on the microstructure and mechanical properties of steel-based composite was investigated. The conclusions drawn from the results can be summarized as follows: 1) Occurrence of DDRX in the IF steel during cold plastic deformation was possible. This phenomenon was due to the presence of SiC microparticles in the IF steel. The four-layer ARB process at room temperature in presence of SiC microparticles provided the stored energy required for DDRX in the IF steel with high SFE. 2) Average grain size of the final sample was 73 nm but the microstructure was relatively inhomogeneous. 3) In the presence of particles, the stored energy and therefore the driving force for CDRX and DDRX increased, and also the large particles acted as nucleation sites through PSN mechanism. 4) After second and third cycles, texture weakening occurred. This is due to occurrence of PSN mechanism. Because of the local 23
constraints imposed by the non-deformable SiC microparticles, these regions undergo different strains and will undergo different crystallographic rotations to the material remote from the particles. 5) First cycle of ARB process had remarkable effect on the dislocation density. The dislocation density increased continually until the dislocation density of the steel-based composite after fourth cycle was about 5.3 times (10.73 × 109 cm-2) higher than that of the initial sample (2.02 × 109 cm-2). The value of dislocation density first rapidly increased, then dwindled, and finally saturated by further ARB cycles. It can be concluded that occurrence of DRV, CDRX, and DDRX concur to keep the dislocation density constant for high strains. 6) After first cycle, a remarkable improve observed in the yield strength value, from 84 MPa to 682 MPa which is almost 8.1 times greater than that of the initial sample. After final cycle, the yield strength value of steel-based composite increased to 1061 MPa. 7) The contribution of grain refinement and precipitation to the improvement in yield strength was maximum (~67%) and minimum (~3.2%), respectively. The predominant method for the improvement of IF steel strength was the grain refining technology. 24
Acknowledgement The corresponding author gratefully acknowledges Mrs. Mahjoobeh Hatef.
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33
Fig. 1. Microstructure of the IF steel at two magnifications.
34
Fig. 2. Schematic illustration of ARB process for fabrication of steel composite.
35
36
37
38
Fig. 3. STEM micrographs from RD–TD plane of samples after (a) first, (b) second, (c) third, and (d) fourth cycles.
39
Fig. 4. Schematic illustration of: (a) boundaries dragged by precipitates and bulging driven by locally concentrated dislocations, and (b) boundaries dragged and bulging driven by low angle boundaries.
40
Fig. 5. The nanoprecipitates in the IF steel microstructure.
41
Fig. 6. Schematic illustration of the deformation path for: (a) regions without particles showing relative straight sliding path, and (b) regions including particles showing longer deformation path due to the rotation of silicon carbide microparticles. 42
43
Fig. 7. {200} pole figures (a) before ARB and after (b) 1, (c) 2, and (d) 3 ARB cycles. 44
45
Fig. 8. Inverse pole figures (a) before ARB and after (b) 1, (c) 2, and (d) 3 ARB cycles.
46
Fig. 9. The h* values for ARB-processed samples calculated by Eq. (1).
47
Fig. 10. The dislocation densities for ARB-processed samples calculated by Eq. (2) based on h* values of Fig. 9.
48
Fig. 11. Engineering stress–strain curves of initial and ARB-processed samples.
49
Fig. 12. Variation of yield strength in ARB-processed samples for different cycles.
50
Fig. 13. Contribution of strengthening mechanisms to the improvement in yield strength of ARB-processed steel-based composite.
51
Tables:
Table 1. Chemical composition of IF steel (in wt.%). C
N
Si
Mn
Cu
Ni
Ti
Fe
0.002
0.004
0.01
0.14
0.01
0.018
0.055
Bal.
52