Dynamic recrystallization in biomedical Co-29Cr-6Mo-0.16N alloy with low stacking fault energy

Author’s Accepted Manuscript Dynamic recrystallization in biomedical Co-29Cr6Mo-0.16N alloy with low stacking fault energy Yunping Li, Yuichiro Koizumi, Akihiko Chiba

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To appear in: Materials Science & Engineering A Received date: 21 March 2016 Revised date: 10 May 2016 Accepted date: 12 May 2016 Cite this article as: Yunping Li, Yuichiro Koizumi and Akihiko Chiba, Dynamic recrystallization in biomedical Co-29Cr-6Mo-0.16N alloy with low stacking fault e n e r g y , Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.05.045 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.

Dynamic recrystallization in biomedical Co-29Cr-6Mo-0.16N alloy with low stacking fault energy Yunping Li1, Yuichiro Koizumi2, Akihiko Chiba2 1.

School of Materials Science and Engineering, Central South University,

Changsha, China, 410083 2.

Institute for Materials Research, Tohoku University, Katahira2-1-1, Aoba-ku,

Sendai, Japan, 980-8577

Abstract Dynamic recrystallization (DRX) behaviors of biomedical Co-29Cr-6Mo-0.16N (CCMN) alloy were analyzed in detail in hot compressive tests. At low temperature ( 1050 oC) and low strain rate ( 1 s-1), DRXed grains preferentially nucleate along pre-existing grain boundaries or 3n boundaries, with a result of necklace microstructure broadening with increasing strain; while at higher temperature and/or higher strain rate, DRX occurs uniformly in matrix resulting in homogenous and texture-free fine grained microstructure. These two kinds of DRXed microstructures formed in CCMN alloy are related the repeated formation of annealing twin during hot deformation, which was determined by both the temperature and the activity of stacking fault formation. It was also found that the 3n boundaries formed at lower strain level will evolve into general high angle grain boundaries (HAGB) with further straining. A schematic model for DRX mechanism of CCMN alloy was proposed with respect to temperature, strain rate, and strain. .

1

Keywords: Co-29Cr-6Mo-0.16N alloy, dynamic recrystallization, annealing twinning, stacking fault, EBSD.

1. Introduction It is well known that working behavior of metallic materials is strongly influenced by stacking fault energy (SFE), composition of alloy and hot working condition. For example, based on the concept that minimum grain size of crystals, dmin during milling of powder is governed by a balance between dislocation generation and the recovery rate arising from dislocation annihilation and recombination, Mohamed [1] pointed out that the normalized minimum grain size dmin/b (b, the Burgers vector) and the normalized SFE /Gb (( , the stacking fault energy and G, the shear modulus) in metals with various SFE follows a logarithmic scale. Following this proposal, a large number of works, aimed at obtaining super-fine grain sized materials, were extended to other processes of metallic materials with low SFE such as Cu-Zn alloys [2-3], Cu-Al alloys [4], and pure copper [5]. SFE plays a great role in deformation process in that deformation of low SFE materials are generally characterized by a large number of localized slipping or stacking faults (SFs). This is in contrast to the relatively homogenous microstructure in high SFE materials, since the cross slipping is inhibited significantly in low SFE materials during deformation [4-7]. In materials with low or extremely low SFE, large formation of SF bands has a high tendency to refine the microstructure through the profuse formations of mechanical twin boundaries or martensitic transformation during deformation or cooling process. Recent researches of Wang on 800H Ni-Fe-Cr-based alloy [8] and especially

2

those of Miura et al. on bicrystal copper via an in-situ tensile test [9] revealed that not the mechanical twinning but annealing twinning plays an important role in the nucleation and growth of DRX. In the bicrystal copper, nucleation of new grains was simply attributed to the 3 annealing twinning growth [9]. Field et al. also conducted a research on the role of annealing twin during static recrystallization of pure copper; the result indicated that annealing twin is not only a product of the recrystallized structure, but also plays an important role in recrystallization process, in which grain growth was heavily dependent upon the annealing twin formation process. A research by Beladi [10] on austenitic Ni-30pct Fe alloy also demonstrated a large number of

3 boundaries in the

DRXed microstructure. However, the dominant DRX mechanism in this alloy was considered to be not the formation of

3 boundaries, but a conventional

nucleation and growth process induced by grain boundary bulging possibly due to the relatively high SFE of this alloy. Yamanaka et al. [11] investigated the DRX behavior of Co-29Cr-6Mo, implying that DRX mechanism is somewhat related to the non-homogenous microstructure due to the extremely low SFE of the alloy. Nevertheless, no clarification on the correlation between annealing twin formation and DRX in Co-29Cr-6Mo alloy was carried out in this research. It has been proved that Co-29Cr-6Mo (hereafter CCM) alloys have extremely low SFE compared to most of the conventional materials [12-14] even at high temperature (approximately -100 mJ.mol

1

and 22 mJ.mol

1

at ambient

temperature and 1000 °C, respectively, if extrapolated from results of Yamanaka [15]). The transition temperature (850°C in equilibrium phase) from

phase to

phase is much higher than that for pure cobalt (417 °C) because the additions of

3

both Cr and Mo reduce the SFE or stabilize the than 1000 °C [16]. It has been found that

phase at temperatures lower

phase is much stabilized with the

slight addition of N [16 17]. The stabilizing effect of

phase was ascribed to the

stable

atoms as

crystalline structure between Co-Cr-N

observed

by

three-dimensional atomic probe [18], indicating that the Cr atom has a much stronger interaction with the interstitial N atom than Co atom, which dynamically delayed the

transition [19]. Nevertheless, SFE of Co-29Co-6Mo-0.16N alloy

(mass%; simply referred to as CCMN alloy hereafter) is still comparable to that of CCM alloy. This is due to the fact that the concentration of N is extremely low, and similar to CCM alloy, a gradual exposed at 800 900

transition in CCMN alloy occurred if

for a long time [20]. Adding N into CCM alloy greatly

helps us to simplify the analysis of DRX process at high temperature, since martensitic transformation or isothermal phase transition scarcely occurred in both deformation above 1000

and in the subsequent rapid cooling process

[21], and a clear clarification the DRX in hot compression would be possible. With CCMN alloy having low SFE even at high temperature, research into the DRX behavior and its mechanism continues to be important. The objective of the present work is to conduct hot compression tests on CCMN alloys, and DRX in relation with 3 annealing twin boundaries will be discussed in details. This has led to a better understanding of DRX through annealing twin boundary formation in low SFE alloy. 2. Experimental materials and procedures The preparation procedure of CCMN alloy has been reported in previous

4

studies [22-23]. We note that N addition was achieved by adding Cr 2N powder into alloy melts. Starting microstructure of the present materials prior to compression has been observed in a previous research [15].The mean grain size was measured to be 52

m (not including

boundaries existed inside the occurred and no thermal more stable

3 boundaries) and a few

3

grains (Fig. 1). No martensitic transformation

phase existed after rapid cooling, indicating the much

phase compared to CCM alloy [14]. Cylindrical specimens, 8 mm

in diameter and 12 mm in height, were machined by electro-discharge machining. To reduce the non-uniformity of microstructure due to friction between the sample surface and jig, the flat ends of the specimen were machined with concentric grooves with a depth of 0.1 mm so that the lubricant in a molten state could flow freely inside the grooves, resulting in a more uniform distribution of microstructure [24]. Compressive tests were carried out in vacuum at temperatures of 1000 °C to 1200 °C with an interval of 50

using a computer-aided Thermecmaster-Z

hot-forging simulator. The selected true strain rates were 0.01, 0.1, 1.0, 10, and 30 s 1. Induction heating on sample was carried out from room temperature to the set temperature at a rate of 5 °C/s. The sample was held at the elevated temperature for 300 s before compression. As soon as the sample was compressed to the final strain level, it was quenched with a mixture of N2 (6 MPa pressure) and He (4 MPa) at a cooling rate of approximately 50 °C/s to room temperature. On the basis of the previous research [24], Moly paste 500 spray (MoS2), and a mixture of Metalglaze347HS, Metalglaze349 and the Moly paste 500 spray were used to reduce the friction coefficient between the specimens

5

and the anvils during the hot forging process. In order to avoid the heat dissipation from sample end surface to the anvil surface at the contact region, mica sheets with a diameter of 20 mm and thickness of approximately 0.2 mm were used as heat insulators closely contacting with the anvil surfaces; they were placed close to the anvil surface. In addition, carbon paper was used as an aid to reduce friction owing to its high heat resistance and high lubricating effect. Crystallographic analysis was conducted by electron backscatter diffraction (EBSD) using an orientation imaging microscope (TexSEM Laboratories, Inc., Provo, UT) attached to the field-emission scanning electron microscope (FESEM). The surface of the specimen for the microstructural observation was prepared by first dry grinding using SiC emery paper. Electrolytic polishing was subsequently conducted in a sulfuric acid and methanol (1:9) solution. TEM observation was carried out in the center of the compression test specimen parallel to the compression axis, using a Topcon EM002B system operated at 200 kV. All TEM foils were cut by electoral discharge machine (EDM) from the center of the hot compressed samples, followed by dimple, ion milling and cleaning in PLASMA CLEANER (model 1020) before loading into the microscope. 3. Experimental Results Two kinds of DRXed microstructures (i,e. necklace microstructure at 1000 o

C and 0.1 s-1 and uniform microstructure at 1100 oC, 30 s-1) after compression

are observed (Fig. 2 a-c). At T=1000 1050 oC, and

= 0.01 1.0 s-1 DRXed

microstructure demonstrates a necklaced microstructure; DRXed grains preferentially occurred around both the pre-existing

3 boundaries and grain

6

boundaries as shown in Fig. 2 (a), and (b). As marked by the red straights in Fig. 2 (a) for the {111} plane traces in the non-DRXed area, most of the deformation bands (marked by dashed straights) occurred within these planes. In addition, both grains II and

as marked in Fig. 2 (b) are pre-existing

3 boundary,

where DRX occurs preferentially along these 3 boundaries, possibly due to the intersections between these boundaries (Fig. 2(a)). TEM observation was conducted in the non-DRXed area of the sample at 1000 oC, 0.1 s-1 and true strain of 0.2 as shown in Fig. 2(d), where these deformation bands along {111} planes were identified to be SF bands; their intersections in different {111} planes are also clearly observed. This implies the localized deformation in terms of SF band formation of present alloy, which is in good agreement with the results of Qu et al. in that the low SFE leads to large density of SF bands [4]. This result also tells us that DRX taking place in the current materials is closely related to the activities of SF bands. At higher temperature and/or higher strain rate, uniformly distributed equi-axial superfine grains after DRX could be observed as shown in Fig. 2 (c) by the inverse polar figure (IPF). The result reveals a texture-free microstructure, and the grain size (including

3n (n=1, 2)

boundaries) was evaluated to be in the range of 1.5 4.2 m (mean grain size of 3.1 m), indicating an extremely uniform DRX inside the grains. Typical necklace microstructure at 1000 oC and strain rate 0.1 s-1 is shown in Fig. 3 at strains of (a ), (a ), (a ), 0.2; (b ), (b ), (b ), 0.4; (c ), (c ), (c ), 0.6; and (d ), (d ), (d ), 1.0, respectively. In order to analyze the DRXed area and non-DRXed area independently, we separated these two areas by EBSD according the following criteria: (1) the new grains are surrounded by high angle 7

grain boundaries (HAGB, misorientation angle above 15 o); (2) the average kernel average misorientation (KAM) value in new refined grains are obviously lower than the mean value of the matrix; (3) since the extensive existence of the 3n boundaries, they were considered to be a kind of HAGB, that is to say, 3n twinning are considered to be new grains if the condition satisfies term (2). At true strain of 0.2, which roughly corresponds to the yield point in the

curve of

current alloy as referred to our previous research [25], traces of deformation bands are clearly observed, and some of them transformed into

3 twinning

plate with relatively low lattice curvature as shown by the red lines for the

3

boundaries (Figs. 3 (a), (a ), (a )). With the increase in strain level, the volume of DRXed

area

increased

gradually

and finally

evolved

into

necklaced

microstructure with a DRXed area fraction of approximately 60% at true strain of 1.0 (Figs. (d), (d ), (d )). It is interesting to observe that most of 3 boundaries appeared in or surrounding the refined grains. Figures 4 show the microstructures at 1100 oC, strain rate 1.0 s-1 and strain levels from 0.2 1.2. A large fraction of 3n boundaries are observed to form in or around the refined grains. Relatively uniform distribution of the refined grains was observed at each strain level and no grain growth was observed with the further compression. This indicates that the DRX take places uniformly inside the matrix. It has to be mentioned that a large number of DRXed grains are surrounded by 3n boundaries, implying that the refining process of CCMN alloy should be closely related to the formation of these

3 boundaries. This is

possibly due to that large formation of SFs via transforming into 3 boundary at high temperature has a high tendency of refining the matrix. This behaves in a

8

way similar to that of the bicrystal pure copper. However, DRX at current condition occur much more homogenously throughout the matrix compared to the pure copper [9], possibly due to the much lower SFE of our alloy. Noted that the DRXed grains with a mean size of approximately 1 m at strain 0.8 were observed to co-exist with the larger grains (but smaller than the starting grains) characterized by low angle grain boundaries (LAGB) inside as indicated by arrow in Fig. 4 (f); the intervals of LAGB was in an order comparable to the size of DRXed grains (Fig. 4(e)). This implies that the DRX proceeded inside the previous DRXed grains again, giving rise to finer grains. Following the result in Figs. 4, we compressed the sample to true strain of 0.4 at same condition as to Fig. 4, a few of super-fine DRXed grains (approximately 500 nm) could be observed under TEM as shown in Fig. 5 (a), in which a large number of SFs formed inside the DRX grains and some grains are surrounded by 3 boundaries. The intersections of SFs inside of the DRX grains are still visible. In the case of 1000 oC at strain rate of 10 s-1 and true strain of 0.4, most of the grains are not recrystallized and we observed a large number of stacking faults, intersecting to each other in the non-DRXed grains (Fig. 5(b)). Area fractions of DRX were calculated at 1000

, and 1100

by EBSD,

and were summarized as a function of true strain level in Fig. 6 (a) and (b), respectively. We did not show the results at 1200 oC since it is very similar to that at 1100 oC. At 1000 oC, before true strain level of 1.0, the maximum fraction was approximately 80% in case of the highest strain rate (30 s -1), while in other conditions, the DRX lower fractions of DRX area were observed. At higher temperature, DRX occurred through the entire matrix, giving rise to uniformly

9

distributed microstructure at strain higher than approximately 0.4 as shown in Fig. 4. Compared to that at lower temperature, area fractions of DRXed microstructure at higher temperature showed higher values (Fig. 6 (b)). It can also be observed that the increasing rate of DRX fraction is highest in the strain range of 0.2-0.6. The DRX fractions as a function of strain rate at true strain of 0.6 are summarized in Fig. 7(a) at various temperatures. It is interesting that the fraction at a strain rate of 0.1 kept the lowest values at lower temperatures (

1050 oC):

the DRXed necklace microstructure as shown in Fig. 2(a) and Fig. 3 seems hard to broaden in this condition. At higher temperature, this tendency is hard to observe due to the full DRX occurring. In the sample deformed at 1000 oC and the lowest strain rate 0.01 s-1 (Fig. 7(b)), larger refined grains are observed when compared to the one at 0.1 s-1 (Fig. 2 (a) and (b)). At low temperature (1000 oC), the higher DRX fraction at strain rate of 0.01 s -1 is possibly due to the longer holding time compared to that at 0.1 s-1. The longer holding time enhanced the transition of 3 annealing twin boundaries from SF bands via diffusion. On the other hand, at high strain rate, the activity of SF formation is enhanced significantly. Together with the temperature rise due to adiabatic heating during such a short time deformation, a transient transition of the SF bands into annealing twin boundary is possible. This has been manifested in previous researches that the temperature rise due to adiabatic heating at strain rate of 10 s-1 can be as high as 50-100 oC at a strain of 0.6 in the case of Co-Cr-Mo alloy [22]. This is because that over 90% of input energy from deformation can transformed into heat, which can enhance the transient nucleation of DRX via

10

the formation of annealing twin boundary. Figures 8 show misorientation angle distributions of microstructure when deformed at (a) 1000 oC, 0.1s-1, and (b) 1100 oC,1.0 s-1, as well as (c) fraction of low angle grain boundaries (LAGB) and fraction of 3n boundaries as a function of true strain level. For both samples obvious peaks of 3n boundaries could be observed at all strain levels. However, in the case of sample at 1100

, 1.0 s-1,

the fractions of 3n boundaries (Fig. 8 (b)) are much higher than that of the other one (Fig. 8 (a)). Correspondingly, fraction of LAGB in the sample at 1100

, 1.0

s-1 are much lower. The fractions of LAGB in both samples decrease gradually with increasing strain (Fig. 9(c)); however, in the sample at 1100

, 1.0 s-1,

fraction of LAGB decreased to 0.1 at a true strain of 0.4, and kept to be roughly a constant, and increased a little with further compression. At the highest strain level of 1.2, the fraction of LABG in the sample at 1000

, 0.1 s-1 is observed to

be still 0.35, indicating the large fraction of residual non-DRXed microstructure. The fractions of

3n boundaries is observed to be low and to vary in an

extremely small range in the sample at 1000

, 0.1 s-1. On the other side,

increasing to a peak value of approximately of 0.49 at a train of 0.6, and the subsequent gradual decrease with strain are observed in the sample at 1100

,

1.0 s-1 (open circles). There are two possible explanations for the results seen in Fig. 8(c) for the sample at 1100

, 1.0 s-1: one explanation is that with an

increasing strain level, the formed

3n boundaries are continuously destroyed

and transformed into general HAGB; the other one might be that the formation of 3n boundaries became more difficult. We observed the microstructure of the

11

sample at 1100 °C, 1.0 s 1, and strain of 0.8 (Fig. 8 (d)) by the grain boundaries map derived from the results of EBSD. The grain boundary angles near

3

twinning boundaries are measured and indicated near the boundaries. Twinning boundaries are observed to loose their coherency, and their angle/axes misorientation relationships began to deviate from the original

3 twinning

misorientation (i.e., 60 deg<111>). The deviations were found to be inhomogeneously distributed along the twinning boundary facets, which strongly support the current hypothesis of twinning boundary evolution into a general HAGB. From the results in Figs. 8(c) and (d), a majority of 3n boundaries lost their unique characteristics and were converted into a general HAGB with straining. Nevertheless, it is considered that both of the abovementioned mechanisms occurred concurrently, since a decreasing grain size inevitably increases the driving force required for twin formation [15]. Furthermore, the slight increase in the LAGB fraction at a high strain level is possibly ascribed to the difficulty in the formation of

3 boundaries and the fact that that the larger

fraction of SFs bands (corresponding to the LAGB) remained stable in the matrix (Fig. 5).

4. Discussion 4.1 Influence of temperature, strain rate and strain on DRX behavior of CCMN alloy At low temperature (lower than 1050 oC) and low strain rate (<1.0 s-1), DRX initiated preferentially near the pre-existing grain boundaries and

3 twin

boundaries along a direction along 35 45 o to the compression direction, leading to necklaced DRX microstructures as shown in Figs. 3. The DRXed 12

microstructure broadened slowly with increase strain, implying that the stress possibly localized in these locations, giving rise to the non-uniform DRX process. This is considered to be related to the fact that at low temperature, large number of SF bands in the matrix could not completely transform into

3 twinning

boundary. This is in contrast to that at high temperature and/or high strain rate, where DRX occurred much more homogenous, possibly due to the easy formation of annealing twin boundary due to high environmental temperature or adiabatic heating. The non-DRXed regions are assumed to represent microstructural features related to the DRXed grains. The microstructure deformed at a temperature of 1000°C, a strain rate of 10 s 1, and a strain level of 0.2 as well as 1100

, strain

rate 30 s-1, strain level of 0.2 are shown in Fig. 11(a) and (b), respectively; Lattice distortions introduced during hot deformation are observed in non-DRXed regions (step size of 0.1 m). Since these lattice misdistortions are mainly due to the large formation of SF bands as mentioned above, we measured the misorientation profiles along the arrows A B, C D, and X Y in the non-DRX grains. Large misorientation angles introduced during deformation, possibly corresponding to grain boundaries in the subsequent DRX process, are sporadically observed. With increasing strain, some of these sub-grain boundaries evolve into 3n twinning boundaries. To evaluate this hypothesis, the point-to-point misorientation profiles were obtained in Figs. 10(c) and (d) along the lines of A B, C D, and X Y, in which the mean value of misorientation through the lines are 2.05 °C, 1.52 °C, and 0.47 °C, respectively. The DRXed grain size of the nonrecrystallized regions, if DRX occurred via the proposed

13

mechanism, would be ,

(1)

where d denotes the grain size after DRX and D, N, and N0 represent the distance of the line, the misorientation peak number, and the misorientation peak number above the mean value of misorientation angle, respectively. The mean sub-grain size of the nonrecrystallized region, if DRX occurred directly through the growth of these sub-grain boundaries, were calculated to be 1.41, 2.76, and 1.84

m respectively, which correspond well with the DRX grain size in the

surrounding area (1.4 2.8

m in Fig. 10 (a), and 1.5 3.0 mm in Fig. 10 (b),

respectively), implying that DRXed grains should grow directly from these SFs bands by the annealing twin formation. We used parameter of power dissipation coefficient

from dynamic

materials model (DDM) of Prasad et al. [26] to explain the stress localization in the mediate strain rate regime (0.1 1.0 s-1) at 1000

, where the parameter

is

given by

and m is the strain rate sensitivity component, and m

where

,

,

log log

(2) ,T

, and T are flow stress, strain rate, strain, and temperature,

respectively. Fig. 9 is the values of

calculated from the deformation curves of

CCMN alloy [25]. It could be observed that at both 1000 oC and 1050 oC, and strain rate 0.1-1.0 s-1,

showed the lowest value. From the DDM, in this case,

there is a high possibility of flow localization in the deformation, since energy

14

input in the deformation was considered do not enhance the microstructure evolution but localized adiabatic heating inside the work pieces, which in turn enhance the DRX process locally- necklaced DRX microstructure. 4.2 DRX mechanism in CCMN alloy Similar to the other research [27-28], Qu et al. [4] conducted a research on the influence of SFE on the microstructure of Cu-Al alloys, indicating that profuse deformation twins, SFs, microscale shear bands (MSB) and their interactions are pronounced in low-SFE Cu-Al alloy because the dislocations can be readily dissociated into partials to form planar arrays of SFs. CCMN alloy at present hot compression condition has comparable SFE to the Cu-Al alloy [4] at ambient temperature. However, in the current study no shear band (SB) or MSB was directly observed in the deformed microstructure, possibly because current test was conducted at high temperature, recovery and/or DRX occurred easily, lowering the stress concentration effectively. This might be proved by the tests conducted at intermediate temperature (700 800

) by T. Odahara [22] on

CCMN alloy, where SBs formed easily, because the stress concentration in the SF bands could not be lowered effectively by recovery or DRX at this temperature. It is interesting to observe that in CCMN alloy at current condition dominant deformation was controlled by SFs formation and their intersections as shown in Fig. 2(d), while deformation twinning was scarcely observed in the deformed matrix in the current conditions. Yamanaka et al. [15] pointed out that DRX of CCM alloy is closely related to the low SFE and that the large local misorientation associated with geometrically necessary (GN) dislocations should be responsible for the nucleation driving

15

force. However, at such low strain, dislocation density, especially the GN dislocation density, was not high enough to form directly the subgrain boundaries through a slipping process. One possible mechanism of DRX in CCMN alloy is related to the large formation of SF bands and their intersections followed by a simultaneous and rapid annealing twin formation process. Since SFE in the present alloy is very low, the stacking faults are assumed to form with extreme ease, with a result that the matrix grains are partitioned into numerous superfine grains through the subsequent formation of finer annealing twins. In fact, Talonen et cl. [29] observed the intersection of stacking faults in metastable austenitic steels at room temperature with SFE varying from 10 to 20 mJ mol 1; their results showed that if SFE of alloy at the deformation temperature is higher than the interfacial energy between mol

1

phase and

phase (approximately 15 mJ

for most alloys), twinning easily occurs. At current deformation

temperatures, CCMN alloy is considered to have an SFE comparable to that of CCM alloy (22 35 mJ mol

1

at temperatures of 1000 °C 1200 °C; if stacking

faults overlap on successive {111} planes, there is a strong possibility of the formation of

3 twin boundaries [30]. In summary, the present DRX process

should be related to the successive 3n twinning boundary formation. In-situ EBSD observation of pure copper after cold deformation followed by heat treatment was carried out by Field et al. [31], the result also indicated that annealing twin boundary formation in materials with low to medium SFE is not only a product of the recrystallized structure, but also plays an important role in the recrystallization process itself. The research by Miura et al. on bicrystal copper in a tensile test at 450 oC also showed that all of the nucleation of DRX

16

was accomplished by means of annealing twin formation along the pre-existing grain boundaries. In the current research at high temperature and/or high strain rate, the nucleation sites are observed to be uniform in the matrix; this is presumed to be due to the lower SFE of CCMN alloy, in that lower SFE leads to higher density of SF and easier formation of annealing twinning, and therefore nucleation sites everywhere in the matrix. From the above-mentioned results, necklaced microstructure after DRX was observed at temperatures

1050

and strain rate

1 s-1(Figs. 2 (a), 3, and

7); this necklace microstructure will broaden with increasing strain level, leading to increase in DRX fraction. At higher temperature, uniform DRX microstructure was observed, implying a strong influence of temperature on the DRX behavior (Fig. 4). In the current temperature range, HAGBs are considered to evolve from 3n boundary, forming rate of which is determined by both temperature and the activity SF bands. At low temperature, for the lower thermal driving force, transition from SF bands into 3n boundaries is more difficult compared to the higher temperature. In this case, DRX process occurred at pre-existing grain boundaries (Fig. 2(a), Fig. 3) or the pre-existing

3 twin boundaries (Fig. 7),

where they gave rise to larger lattice misorientation than that intersection between SFs bands in the matrix-the local DRX process. Higher strain rate increased the number of SFs from the results of Murr et al. and Hecker et al. in 304 stainless steel [32-33], and subsequently provides higher driving force for transition from SF bands into 3n boundaries for the high density of SF bands. Similarly, our results in Fig. 5 (a) and (d) show that with the increase in strain rate (10 s-1, and 30

-1

), higher fraction of DRX was observed compared to that at

17

lower strain rate (0.1 s-1, and 1.0 -1). On the other side, we also observed a slight increase in DRX fraction if the strain rate is in an extremely low level (0.01 s-1 in Fig. 6 (a), (d)). This result is considered, as mentioned previously, to be related to that long deformation time greatly enhanced the transformation of SFs into 3n boundaries. Research by Odahara. [22] greatly helps us to explain this phenomena: in the deformed CCMN alloy at ambient temperature, static recrystallization initiated if the deformed sample was kept at 1000

for a short

time. The growth into HAGB (dominant 3n boundaries) was closely related to the intersection of SFs, and if the compressed sample was kept at higher temperature (1100

), the growth of the SF bands into

3 boundaries is a

transient process [22]. This is also the reason why very uniform DRXed microstructure was obtained at higher temperature. According to the above discussion, a model for the DRX mechanism of CCMN alloy was proposed as shown in Fig. 11. At low temperature, the stress concentration near the intersections between grain boundaries or large 3 twining boundaries with the SFs provide the preferential nucleation sites for DRX process, and the DRXed grain size is roughly equal to the mean bands width of the SFs; while at high temperature, due to the high thermal driving forces for the growth of LAGB into HAGB, the intersection between SFs with original grain boundaries and

3

boundaries just play a minor role, and DRX mechanism evolve into a growth of 3 twinning boundaries from the intersections of SF bands, which proceed in a cyclic way with a result of

extremely fine grains at high strain level (Fig. 5).

5. Conclusions In the current study, dynamic recrystallization (DRX) behaviors were analyzed in

18

detail in compression tests of biomedical Co-29Cr-6Mo-0.16N (CCMN) alloy at temperature of 1000 to 1200 oC, and strain rate of 0.01 s-1 to 30 s-1. The results summarized are as follows: At low temperatures and low strain rate, DRX grains preferentially nucleate along original grain boundaries or

3 boundaries, with a result of necklace

microstructure broadening with increasing strain, while at high temperature or high strain rate, DRX occurs uniformly inside the matrix with a result of homogenous texture-free fine grained microstructure. The DRX mechanism was explained by a schematic model as a function of temperature, strain rate, and strain. Acknowledgements This research was partially supported by a Cooperation of Innovative Technology and Advanced Research in Evolutional Area from Ministry of Education, Culture, Sports, Science and Technology of Japan. References [1] Mohamed FA, A dislocation model for the minimum grain size obtainable by milling, Acta Mater. 51 (2003) 4107-4119. [2] Zhao YH, Zhu YT, Liao XZ, Horita Z, Langdon TG, Influence of stacking fault energy on the minimum grain size achieved in severe plastic deformation, Mater Sci Eng A, 463(2007) 22 26. [3] Balogh L, Unga T, Zhao Y, Zhu YT, Horita Z, Xu C, Langdon TG, Influence of stacking-fault energy on microstructural characteristics of ultrafine-grain copper and copper zinc alloys, Acta Mater, 56(2008) 809 820. [4] Qu S, An XH, Yang HJ, Huang CX, Yang G, Zang QS, Wang ZG, Wu SD,

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Zhang ZF, Microstructural evolution and mechanical properties of Cu Al alloys subjected to equal channel angular pressing, Acta Mater, 57(2009) 1586 1601 [5] Torre FD, Lapovok R, Sandlin J, Thomson PF, Davies CHJ, Pereloma EV. Microstructures and properties of copper processed by equal channel angular extrusion for 1 16 passes, Acta Mater, 52(2004), 4819-32. [6] Belyakov A, Sakai T, Miura H, Kaibyshev R, and Tsuzaki K, Continuous recrystallization in austenitic stainless steel after large strain deformation, Acta Mater, 50(2002), 1547-57 [7] Sommitsch C, and Mitter W, On modelling of dynamic recrystallisation of fcc materials with low stacking fault energy, Acta Mater, 54 (2006) 357 375 [8] Wang X, Brunger E, and Gottstein G: The role of twinning during dynamic recrystallization in alloy 800H, Scripta Mater, 46(2002), 875 80. [9] Miura H, Sakai T, Mogawa R, and Gottstein G, Nucleation of dynamic recrystallization at grain boundaries in copper bicrystals, Scripta Mater, 51(2004), 671 675. [10] Beladi H, Cizek P, and Hodgson P, Dynamic Recrystallization of Austenite in Ni-30 Pct Fe Model Alloy: Microstructure and Texture Evolution, Metall. Mater. Trans. A, 40A(2009), 1175 89. [11] Yamanaka K, Mori M, Chiba A, Origin of Significant Grain Refinement in Co-Cr-Mo Alloys Without Severe Plastic Deformation, Metallurgical and Materials Transactions A, 43(2012) 4875-7887. [12] Saito Y, Tsuji N, Utsunomiya H, Sakai T, and Hong RG, Ultra-fine grained bulk aluminum produced by accumulative roll-bonding (ARB) process, Scripta Mater, 39(1998), 1221 27.

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[13] Saito Y, Utsunomiya H, Tsuji N, and Sakai T, Novel ultra-high straining process for bulk materials

development of the accumulative roll-bonding (ARB)

process, Acta Mater, 47(1999), 579 83. [14] Xing J, Yang X, Miura H, and Sakai T, Mechanical Properties of Magnesium Alloy AZ31 after Severe Plastic Deformation, Mater. Trans., 49(2008), 69 75. [15] Yamanaka K, Mori M, Kurosu S, Matsumoto H, and Chiba A, Ultrafine Grain Refinement

of

Biomedical

Co-29Cr-6Mo

Alloy

during

Conventional

Hot-Compression Deformation, Metall. Mater. Trans. A, 40A (2009), 1980-94. [16] Lee S, Nomura N, and Chiba A, Significant Improvement in Mechanical Properties of Biomedical Co-Cr-Mo Alloys with Combination of N Addition and Cr-Enrichment, Mater. Trans. 49(2008), 260-64 [17] Shi LQ, Northwood DO, and Cao ZW, Alloy design and microstructure of a biomedical Co-Cr alloy, J. Mater. Sci. 28(1993) 1312 16. [18] Li Y, Yu JS, Kurosu S, Koizumi Y, Matsumoto H, Chiba A, Role of N in stabilizing the phase of Co-Cr-Mo alloy, Mater Chem Phys, 133 (2012), 29-32.

[19] Kilner T, Dempsey AJ, Pilliar RM, Weatherly GC, The effects of nitrogen additions to a cobalt-chromium surgical implant alloy, J Mater Sci., 22(1987), 565-574. [20] Kurosu S, Matsumoto H, Chiba A, Grain refinement of biomedical Co 27Cr 5Mo 0.16N alloy by reverse transformation, Mater Letters, 64(2010), 49-52. [21] Chiba A, Biomedical Co-Cr-Mo alloys, Materia Japan, 46(2007), 194 97. [22] Odahara T, Development of Co-Cr-Mo alloy with high concentration of Cr, Doctor dissertation, 2010, Tohoku University. [23] Chiba A, Lee HS, and Matsumoto H, Construction of processing map for

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biomedical Co 28Cr 6Mo 0.16N alloy by studying its hot deformation behavior using compression tests, Mater Sci Eng. A, 513-514(2009), 286-293. [24] Li Y, Onodera E, Matsumoto H, and Chiba A, Correcting the Stress-Strain Curve in Hot Compression Process to High Strain Level, Metall. Mater. Trans. A, 40A (2009), 982-90. [25] Li Y, Onodera E, Matsumoto H, and Chiba A, Construction of processing map for biomedical Co-Cr-Mo alloy, Journal of the Japan Society for Technology of Plasticity, 51(2010), 221-26. [26] Prasad YVRK, Gegel HL, Doraivelu SM, Malas JC, and Morgan JT, Modeling of dynamic material behavior in hot deformation: Forging of Ti-6242, Metall. Trans. A, 15A (1984), 1883 92. [27] Wang XY, Li YP, Hou YH, Koizumi Y, Chiba A, Effects of surface friction treatment on the in vitro release of constituent metals from the biomedical Co-29Cr-6Mo-0.16N alloy, Mater Sci Eng, 64C(2016), 260-268. [28] Hou YH, Li YP, Zhang C, Chiba A, Ex-situ observation on the dissolution behavior of Ni-16Cr-15Mo and Ni-30Co-16Cr-15Mo alloys in aqueous hydrofluoric acid solution, Corros Sci, 90(2015), 133-139. [29] Talonen J, Hänninen H, Formation of shear bands and strain-induced martensite during plastic deformation of metastable austenitic stainless steels Acta Mater, 55(2007), 6108-18. [30] Mahajan S, Pande CS, Imam MA, Rath BB, Formation of annealing twins in f.c.c. crystals, Acta Mater, 45(1997) 2633-2642. [31] Field DP, Bradford LT, Nowell MM, Lillo TM, The role of annealing twins during recrystallization of Cu, Acta Mater, 55 (2007) 4233 4241.

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[32] Murr LE, Staudhammer KP, Hecker SS, Effects of Strain State and Strain Rate on Deformation-Induced Transformation in 304 Stainless Steel: Part II. Microstructural Study, Metall. Trans. A, 13(1982), 627-35. [33] Hecker SS, Stout MG, Staudhammer KP, Smith JL, Effects of Strain State and Strain Rate on Deformation-Induced Transformation in 304 Stainless Steel: Part I. Magnetic Measurements and Mechanical Behavior, Metall. Trans. A, 13(1982), 619-25.

Fig. 1 (a) Initial microstructure of CCMN alloy after 1100

300 s annealing

followed by rapid cooling, and (b) the corresponding XRD pattern. In (a) the black line represents the HAGB, and red line represents the

3 annealing

twinning boundaries.

Fig. 2 (a), IQ map and (b) IPF map of CCMN alloy at 1000 s-1, and true strain of 0.6. (a) HABG and

, strain rate of 0.1

3 boundaries are marked by white

lines and red lines, respectively; (b) HABG and S3 boundaries are marked by black lines and red lines, respectively. (111) plane traces are marked with red straights along different directions in (a), (c) IPF map of CCMN alloy at 1100

,

30 s-1, to strain of 0.6(d) TEM microstructure of CCMN deformed at 1000

,

strain rate of 0.1 s-1, and true strain of 0.2.

Fig. 3 Microstructure of CCMN alloy at 1000

and strain rate of 0.1 s-1 and

PF maps

23

non-

Fig. 4 Microstructure of CCMN alloy at 1100

and strain rate of 1.0 s -1 and

true strain levels of (a) 0.2, (b), 0.4, (c), 0.8, (d) 1.2; IQ map with HAGB marked with HAGB (black lines) and

3 twinning boundaries (red line), S9 twinning

boundaries (blue lines). Fig.5 TEM micrographs of the Co-29Cr-6Mo-0.12N alloy after hot forging at strain rate of 10 s-1 (true strain:0.4: (a) at 1100 oC and (b) 1000 oC.

Fig. 6 DRX area fraction as a function of true strain, (a) 1000

; (b) 1100.

Fig. 7 (a) DRX area fraction as a function of strain rate, (b) IPF map of CCMN alloy at 1000

, strain rate of 0.01 s -1, and true strain of 0.6. HABG and S3

twinning boundaries are marked by black lines and red lines, respectively.

Fig. 8 Misorientation angle distributions deformed at (a) 1000 1100

, 0.1 s-1, and (b)

, 1.0 s-1, at various strain levels, and (c) Number fraction of LAGB and

Sigma 3 boundaries as a function of true strain levels, (d) grain boundary map after deformed at 1100

, 1.0 s-1, to strain of 0.8.

Fig. 9 Power dissipation coefficient h as a function of true strain rate.

Fig. 10 (a) Microstructure of sample deformed at 1000 strain of 0.2, (b) Microstructure of sample deformed at 1100

and 10 s-1 with a and 30 s-1 with a

strain of 0.2, (c) point to point misorientation profile along the arrows of A-B (black solid line), C-D (red dashed line), and (d) point to point misorientation profile along the arrows of X-Y (black solid line), respectively. Fig. 11 DRX model developed for CCMN alloy.

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