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Age-hardening mechanism for nanocrystalline Ni–P alloys synthesized by electrodeposition
SCT-19416; No of Pages 7 Surface & Coatings Technology xxx (2014) xxx–xxx
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Age-hardening mechanism for nanocrystalline Ni–P alloys synthesized by electrodeposition Yosuke Kasazaki, Hiroshi Fujiwara, Hiroyuki Miyamoto ⁎ Department of Mechanical Engineering, Doshisha University, 1-3 Miyakodani, Tatara, Kyotanabe, Kyoto 610-0321, Japan
a r t i c l e
i n f o
Article history: Received 17 February 2014 Accepted in revised form 12 May 2014 Available online xxxx Keywords: Nanocrystal Electrodeposition Ni–P alloy Precipitation hardening
a b s t r a c t Strengthening mechanism of age-hardenable electrodeposited Ni–P alloy was investigated focusing on the role of Ni3P precipitates in nanocrystalline structures. Specimens of P content ranging from 3.0 to 12.5 wt.% were synthesized by electrodeposition followed by aging treatment at a temperature between 473 and 773 K. The specimens became harder with higher P content after the electrodeposition and aging treatment. As-electrodeposited structures comprised only Ni phase, and Ni3P precipitated by the subsequent aging treatment. Vickers hardness of electrodeposits showed the maximum after aging at 673 K, and was higher than that predicted by the rule of mixture of Ni and Ni3P phases. Distribution of parent Ni grain size and Ni3P precipitate size was examined in detail by field-emission type transmission electron microscopy. The maximum hardness was obtained when the grain size of Ni matrix and Ni3P particle size were comparable. The hard interphase boundary to the dislocation slip is included as a possible mechanism which is responsible for the hardening by the aging treatment rather than the classical precipitation hardening. © 2014 Elsevier B.V. All rights reserved.
1. Introduction Electrodeposited nickel–phosphorus coatings have long been used for applications where a hard coating is required, as they are remarkably hard compared to pure nickel coatings [1–5]. Ni–P alloys become dramatically harder when subjected to an aging treatment following electrodeposition [6–12]. Phosphorous has an extremely low solubility limit in nickel, and the alloy is an amorphous, supersaturated solid solution after electrodeposition [3,6,13–16]. Subsequent aging causes particles such as Ni3 P to precipitate, and the alloy simultaneously assumes a nanocrystalline structure [13–15]. The strengthening effect has often been explained simply as precipitation hardening, similar to that which occurs in materials with much larger grains [7–9], although a few other mechanism was reported [11]. However, very little detailed research has been carried out into the strengthening mechanism [7–10]. We have previously investigated the Ni and Ni 3P grain size distributions following aging of Ni–P alloys produced by electrodeposition, and concluded that the strengthening was not due to precipitation hardening, but instead due to the interface between the Ni matrix and the Ni3P precipitates, which have different crystal structures [6]. However, some doubts remain about this conclusion because only specimens with a
⁎ Corresponding author. E-mail address: [email protected] (H. Miyamoto).
single P content were used for dark-field transmission electron microscopy observations of Ni3P precipitates. Therefore, it is necessary to examine the strengthening mechanism in more detail using specimens with a wider range of precipitate volume fractions. There are currently several proposals for the deformation mechanism in nanocrystalline materials, such as grain boundary sliding and the formation/disappearance of dislocations at grain boundaries. For this reason, the strengthening mechanism in Ni–P composites is an extremely interesting subject of research, because it is likely to provide important insights into the deformation mechanism, and thus greatly contribute to the future development of nanomaterials. In the present study, the strengthening mechanism in Ni–P alloys was examined by forming Ni–P films by electrodeposition, and then producing ultrafine, two-phase Ni–Ni3P alloys by aging under various conditions. 2. Experimental methods In this study, Ni–P alloy specimens were produced by electrodeposition. Table 1 shows the bath composition and electrodeposition conditions used. The P content in the alloys was changed by varying the amount of H3PO3 added to the bath, and H3PO4 was also added to adjust the bath pH to 1.5. The anode and cathode substrates used during electrodeposition were Ni (99.9%) and stainless steel (SUS 304) having a size of 20 × 30 mm, respectively, and they were ultrasonically cleaned in acetone prior to use. The bath was stirred constantly during electrodeposition using a magnetic stirrer, in order to homogenize its composition and temperature, and to withdraw
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Ni(SO3NH2)2–4H2O [g/ℓ] NiCl2–6H2O [g/ℓ] H3BO3 [g/ℓ] H3PO3 [g/ℓ] Plating temperature [K] Current density [mA/cm2] Type of current Plating time [h]
P content [mass%]
Table 1 Bath composition and deposition conditions. 470 15 10 1, 3, 5, 7, 10, 15 323 200 D.C. 2
10
5
0 0 the gas generated by ionization away from the cathode surface to prevent pitting. Deposition was carried out using a pulsed power supply (HCP-301H, Hokuto Denko Corporation). The electrolyte used was a mixed liquid comprising 40% acetic acid, 30% phosphoric acid, 20% nitric acid, and 10% distilled water. As shown in the thermal equilibrium diagram for Ni–P alloy in Fig. 1 [17], the amount of P in solid solution in Ni is greatest (0.32 at.%) at around 900 °C, and P hardly forms a solid solution at room temperature. Therefore, the as-deposited film is a supersaturated solid solution, and aging causes precipitation of Ni3P. In this experiment, the aging treatment involved heating the specimens under a vacuum for 40 min at temperatures of 473 to 773 K. X-ray diffraction (XRD; SmartLab, Rigaku Corporation) analysis was performed in order to determine the crystal structure and the grain size in the specimens. The diffraction conditions were a voltage of 40 kV, a current of 200 mA using CuKα, a scan speed of 4°/min, and an angular range of 30 to 80°. Hardness measurements were carried out using a micro-Vickers hardness tester (HMV-2200, Shimadzu Corporation) under constant load of 980.7 mN. Microstructural observations were performed using transmission electron microscopy (TEM; JEM-2100F, JEOL Ltd.) at an accelerating voltage of 200 kV, and local compositional measurements were carried out using energy-dispersive X-ray spectroscopy (EDS) in the TEM system. The P content in the specimens was determined using inductively coupled plasma optical emission spectrometry (ICP-OES).
3. Experimental results Fig. 2 shows the relationship between the amounts of H3PO3 in the bath and the P concentration after electrodeposition, as determined by
L
1200
31
800 0.32
880
19
(Ni)
Ni2P
400
Ni3P Ni5P2 Ni12P5
600
200
850 825 770
NiP3
870
G
700 450
NiP2
23.5
1000
Ni5P1 Ni P 1.22
Temperature [ ]
1600 1455 1400
(P)
0 0
Ni
10
20
30
40
50
60
70
80
at%P
Fig. 1. Thermal equilibrium phase diagram for Ni–P [17].
90 100
P
5
10
15
20
Fig. 2. Dependence of phosphorous content in as-deposited Ni–P on amount of H3PO3 added.
ICP-OES. The P content is seen to increase almost linearly with the amount of H3PO3 as was reported elsewhere [7]. Fig. 3 shows XRD patterns for as-deposited and aged specimens with different P contents. When the P content is low, the peaks are sharp before and after the heat treatment, showing that the specimens are polycrystalline. However, with increasing P content, peak broadening is observed for the asdeposited specimens, indicating the presence of an amorphous state. The heat treatment causes crystallization to take place, as evidenced by sharpening of the peaks. The temperature at which this occurs tends to increase with P content. For the specimen with 0.8 wt.% P, only peaks associated with Ni were identified. However, for P contents of 5.0, 6.1, 7.7, and 12.5 wt.%, Ni3P peaks were apparent for specimens aged at 673 K. Fig. 4 shows the effect of the P content on the grain size in the Ni parent phase, as determined from the XRD results using the Scherrer formula. For the as-deposited specimens and those heat treated at 473 K, the grain size decreased with increasing P content. However, for the specimen heat treated at 673 K, the grain size was independent of P content. Fig. 5 shows the relationship between the aging temperature and the micro-Vickers hardness. For 0.8%P, the hardness increased monotonically with aging temperature whereas the hardness of higher P contents increased with aging temperature, with the maximum hardness being obtained at 673 K. Figs. 6 to 8 show bright-field, dark-field TEM images, and selected area electron diffraction patterns (SADP) for specimens with a P content of 5.0 mass% after electrodeposition, and aging treatment at 473 K and 673 K as an example. In the diffraction pattern in Fig. 6(b) for the asdeposited specimen, only Debye rings associated with Ni were identified. The dark-field image in Fig. 6(c) was taken using the region of the Debye ring indicated by the circle in the diffraction pattern. From dark-field observations, the average Ni grain size was determined to be 5 nm. For the specimen heat treated at 473 and 573 K, the SADP again contained only Debye rings associated with Ni, and the average grain size determined from the dark-field images was 5 and 9 nm, respectively. For specimens heat treated at 673 and 773 K, diffraction spots associated with Ni and Ni3P were identified, and the average grain size for the two phases was 25 and 23 nm, respectively, for the specimen heat treated at 673 K, and 85 and 74 nm, for the specimen heat treated at 773 K. Fig. 9 shows a compilation of TEM results for specimens with different P contents heat treated at 673 K, which is the temperature that yielded the maximum hardness. For a P content of 0.8 wt.%, no Ni3P was observed, and the average Ni grain size was 35 nm. For a P content of 3.0 wt.% and above, Ni and Ni3P were observed, and the average grain size of both Ni and Ni 3P was about 25 nm, which is comparable to that calculated using the Scherrer formula.
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(200) (111) (200)
(111)
(220)
(a) 0.8 mass%
(220)
(b) 3.0 mass% (220)
(111) (200)
(220)
(c) 5.0 mass%
(111)
(200)
(d) 6.1 mass%
(111) (200)
(220)
(e) 7.7 mass%
(f) 12.5 mass%
Fig. 3. XRD patterns for as-deposited and aged specimens with P contents of (a) 0.8 mass%, (b) 3.0 mass%, (c) 5.0 mass%, (d) 6.1 mass%, and (e) 12.5 mass%.
4. Discussion 4.1. Comparison with rule of mixtures Assuming that all P is precipitated as Ni3P after the aging treatment, the volume fraction of Ni3P will be 6.0 to 95.4 vol.% depending on the P
content. Therefore, when considering the strengthening mechanism for the Ni–Ni3P alloy, the rule of mixtures for Ni and Ni3P cannot be ignored. Fig. 10 shows the values obtained by the experiments and those predicted by the rule of mixtures for specimens aged at a temperature of 673 K. The Ni grain size was taken to be the average of the grain sizes estimated using the Scherrer formula and measured directly from the dark-field
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1200
40
0.8 P (mass%) 3.0 5.0 7.7
30
Vickers hardness [HV]
Grain size [nm]
35 25 20 15 10 5 0 0
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P contents [mass%]
600 400 200 0 200
(a) As-deposit
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Aging temperature [K]
40 35
Grain size [nm]
800
Fig. 5. Dependence of micro-Vickers hardness on aging temperature.
30 25 20 15 10 5 0 0
2
4
6
8
10
12
14
P contents [mass%]
Vickers hardness values for Ni and Ni3P with a grain size of 25 nm were determined to be 530 and 950 HV, respectively. The predicted Vickers hardness for different Ni3P volume fractions is indicated by the dashed line in Fig. 10. It can be seen that the measured values are larger than the predicted values for all compositions. Therefore, it is possible that some additional strengthening mechanism occurs in the Ni– Ni3P alloy.
(b) 473 K 4.2. Grain size distribution in Ni–Ni3P alloy
40
Grain size [nm]
35
Age hardening of Ni–P alloys has often been explained as classic precipitation hardening due to the presence of Ni3P, similar to that which
30 25 20 15
(a)
10
(b)
5 0 0
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P contents [mass%]
(c) 673 K Fig. 4. Dependence of grain size on P content for (a) as-deposited specimens, (b) specimens heat treated at 473 K, and (c) specimens heat treated at 673 K.
100nm
(c) TEM images. The average grain sizes for specimens with a P content of 0.8, 3.0, 5.0, 6.1, 7.7, and 12.5 wt.% were 35, 24, 25, 26, 26, and 24 nm, respectively. Since the grain size of the specimen with the 0.8 wt.% P content was exceptionally large, the hardness value was excluded when comparing with the rule of mixtures. The Vickers hardness for single phase Ni and Ni3P with a grain size of 25 nm was calculated using the method reported in Ref. 11, assuming that the Hall–Petch law was valid. The Hall–Petch coefficient k is reported to be proportional to the square root of the shear modulus [8]. Therefore, k for Ni3P was determined from the known value for Ni, using the ratio of the shear moduli of the two materials. Using a value of 64 GPa for the shear modulus for Ni3P [18], the relationship kNi3P = 0.92kNi was obtained, and the
100nm Fig. 6. TEM results for as-deposited specimens with a P content of 5.0 mass%: (a) brightfield image, (b) diffraction pattern, and (c) dark-field image of Ni grains obtained from the region indicated by the circle in the diffraction pattern.
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(a)
(b)
5
(a)
(b) A B
100nm
(c)
100nm
(c)
100nm
(d)
100nm
100nm
Fig. 7. TEM results for specimens with a P content of 5.0 mass% heat treated at 473 K: (a) bright-field image, (b) diffraction pattern, and (c) dark-field image of Ni grains obtained from the region indicated by the circle in the diffraction pattern.
Fig. 8. TEM results for specimens with a P content of 5.0 mass% heat treated at 673 K: (a) bright-field image, (b) diffraction pattern, (c) dark-field image of Ni grains obtained from region A in the diffraction pattern, and (d) dark-field image of Ni3P grains obtained from region B in the diffraction pattern.
occurs in materials with much larger grains. However, for this to be the case, the Ni3P grains would need to be considerably smaller than the Ni grains, and to exist within their interior. Otherwise, the Ni3 P grains could not impede dislocation slip. Therefore, the grain size distribution for Ni and Ni3P was determined from TEM dark-field images of specimens that had been heat treated at 673 K, which is the temperature that produced the maximum hardness. Fig. 11 shows the Ni and Ni3P grain size distributions for specimens with P contents of 3.0, 5.0, and 6.1 wt.%. Although a wide distribution is obtained for all specimens, there is no significant difference between the average and median values for Ni and Ni3 P. Therefore, although there may be very small second-phase particles partially dispersed inside the Ni grains, it is considered unlikely that precipitation hardening is the dominant strengthening mechanism. One possibility is that strengthening occurs due to the heterophase interface between the cubic Ni grains and the tetragonal Ni 3 P precipitates [19]. The extra increase of hardness over the rule-of-mixture prediction, which already takes into account the strengthening effects of nanocrystalline structure, must be related to the Ni–Ni3 P interphase boundaries. It is rational to consider that the interface between dissimilar phases (cubic–tetragonal) is more effective than homophase grain boundaries in impeding slip transmission across the interface [19]. The Burgers vectors for dislocations on either side of the interphase interface are different, so the interface may strongly suppress dislocation movement, in contrast to the case for a single-phase grain boundary. The higher density of “geometrically necessary dislocations” at the interface as a result of the inherently greater strain incompatibility between adjacent grains of dissimilar materials can also make the interface harder [19,20]. In Fig. 10, it can be seen that the largest differences between the measured and predicted hardness values are obtained when the volume fractions of Ni and Ni3P are roughly the same. In such a situation, the heterophase interface makes up the largest fraction of the total interface area, and would
therefore be expected to have the largest effect on dislocation movement. Both the strengthening mechanism argued by the present authors and by Chang et al. [11] exclude dislocation pinning by Ni3P precipitation as the primary role. However, there is a clear difference. In Chang's case, the change from non-equilibrium to equilibrium grain boundaries during the annealing is the cause of the strengthening. In as-electrodeposited state, grain boundaries are at non-equilibrium state with extrinsic grain boundary dislocations and trapped extra dislocations. These non-equilibrium grain boundaries can operate as dislocation sources in deformation, and as a result they become transplant for crossing dislocations. Thus, the relaxation at non-equilibrium grain boundaries on annealing makes it more difficult for them to dislocate. The so-called “hardening by annealing” phenomenon has been reported for metals by electrodeposition [21,22] and severe plastic deformation [23–25]. However, it is important to note that the degree of strengthening by annealing in electrodeposited and severely deformed pure metals was no more than 10%, and was much smaller than that in the present study and that reported by Chang et al. [11]. Therefore, the effect of structural change from non-equilibrium to equilibrium state of grain boundaries on the strengthening should be limited in the Ni– P system though it is not negligible. Thus, it is rational to consider that P atoms either as solute or second phase plays a primary role in the strengthening upon the annealing. Although it has been argued that all P is precipitated as Ni3P, it is also possible that some P segregates at grain boundaries [16]. This may also influence the material strength, because grain boundaries are involved in the deformation mechanism for nanocrystalline materials. Fig. 12 shows a TEM–EDS analysis of a specimen with a P content of 5 wt.%, that was aged at a temperature of 673 K. The Ni and P contents were measured along the line profile indicated in red. It can be seen that the P signal is strong within regions with a width of about 30 nm, which corresponds to the size of Ni3 P precipitates.
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P=0.8 mass%
P=3.0 mass%
P=6.1 mass%
P=7.7 mass%
P=12.5 mass%
(a) 100nm
100nm
(b)
A
B
A B
100nm
100nm
100nm
B
A B
B
A
A
(c) 100nm
100nm
100nm
100nm
100nm
100nm
100nm
100nm
100nm
100nm
(d)
Fig. 9. TEM results for specimens heat treated at 673 K: (a) bright-field images, (b) diffraction patterns, (c) dark-field images of Ni grains obtained from region A in the diffraction patterns, and (d) dark-field images of Ni3P grains obtained from region B in the diffraction patterns.
Also, the sizes of the regions with strong P and Ni signals are comparable,. However, there is no evidence for localized P signal at grain boundaries. This indicates that the amount of segregation of P to grain boundaries is small, and its effect on strengthening can be considered low.
P content [mass %] 0
5
10
Vickers hardness [HV]
1200 1000 800 600 400
Ni-Ni3P experimental value Rule of mixture of Ni-Ni3P (d=25µm)
200
5. Summary To clarify the age hardening mechanism in Ni–P alloys, specimens with different P contents formed by electrodeposition were subjected to aging treatments at different temperatures. The hardness was found to increase with aging temperature, and the amount of Ni3P precipitated increased with P content. Regardless of the P content, the maximum hardness was obtained for specimens aged at 673 K, and for these specimens, the Ni and Ni3 P grain size distributions were evaluated by X-ray diffraction and electron microscopy. The results showed that the grain size distribution was similar for Ni and Ni3P. This makes it unlikely that precipitation hardening, in which dislocation movement is hindered by small particles within grains, plays a significant role. In addition, the hardness after aging was higher than that predicted using the rule of mixtures for all Ni3P volume fractions. The largest differences occurred when the volume fractions of the two phases were almost the same. This suggests the possibility that the interface between the Ni and Ni 3P phases plays a large role in the strengthening mechanism.
0 0
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40
60
80
100
Acknowledgments
Volume fraction of Ni3P [vol%] Fig. 10. Measured Vickers hardness and that predicted by the rule of mixtures.
The authors gratefully acknowledge the financial support of a Grant-in-Aid for Scientific Research on Innovative Areas “Bulk nano
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(a) 0.25 N=100
Frequency
0.2 0.15
Ni Ni3P
0.1 0.05 0
Grain Size [nm]
(b) 0.25 N=100
Frequency
0.2
Fig. 12. TEM image and EDS line profiles for Ni and P in a specimen aged at 673 K. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
0.15 0.1
Ni Ni3P
0.05 0
Grain Size [nm]
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(c)
[8]
0.25 N=100
Frequency
0.2 0.15
Ni Ni3P
0.1
[9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
0.05 0
Grain Size [nm] Fig. 11. Ni and Ni3P grain size distributions in specimens aged at 673 K, with P contents of (a) 3.0 mass%, (b) 5.0 mass%, and (c) 6.1 mass%.
[19] [20] [21] [22] [23] [24]
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Age-hardening mechanism for nanocrystalline Ni–P alloys synthesized by electrodeposition Yosuke Kasazaki, Hiroshi Fujiwara, Hiroyuki Miyamoto ⁎ Department of Mechanical Engineering, Doshisha University, 1-3 Miyakodani, Tatara, Kyotanabe, Kyoto 610-0321, Japan
a r t i c l e
i n f o
Article history: Received 17 February 2014 Accepted in revised form 12 May 2014 Available online xxxx Keywords: Nanocrystal Electrodeposition Ni–P alloy Precipitation hardening
a b s t r a c t Strengthening mechanism of age-hardenable electrodeposited Ni–P alloy was investigated focusing on the role of Ni3P precipitates in nanocrystalline structures. Specimens of P content ranging from 3.0 to 12.5 wt.% were synthesized by electrodeposition followed by aging treatment at a temperature between 473 and 773 K. The specimens became harder with higher P content after the electrodeposition and aging treatment. As-electrodeposited structures comprised only Ni phase, and Ni3P precipitated by the subsequent aging treatment. Vickers hardness of electrodeposits showed the maximum after aging at 673 K, and was higher than that predicted by the rule of mixture of Ni and Ni3P phases. Distribution of parent Ni grain size and Ni3P precipitate size was examined in detail by field-emission type transmission electron microscopy. The maximum hardness was obtained when the grain size of Ni matrix and Ni3P particle size were comparable. The hard interphase boundary to the dislocation slip is included as a possible mechanism which is responsible for the hardening by the aging treatment rather than the classical precipitation hardening. © 2014 Elsevier B.V. All rights reserved.
1. Introduction Electrodeposited nickel–phosphorus coatings have long been used for applications where a hard coating is required, as they are remarkably hard compared to pure nickel coatings [1–5]. Ni–P alloys become dramatically harder when subjected to an aging treatment following electrodeposition [6–12]. Phosphorous has an extremely low solubility limit in nickel, and the alloy is an amorphous, supersaturated solid solution after electrodeposition [3,6,13–16]. Subsequent aging causes particles such as Ni3 P to precipitate, and the alloy simultaneously assumes a nanocrystalline structure [13–15]. The strengthening effect has often been explained simply as precipitation hardening, similar to that which occurs in materials with much larger grains [7–9], although a few other mechanism was reported [11]. However, very little detailed research has been carried out into the strengthening mechanism [7–10]. We have previously investigated the Ni and Ni 3P grain size distributions following aging of Ni–P alloys produced by electrodeposition, and concluded that the strengthening was not due to precipitation hardening, but instead due to the interface between the Ni matrix and the Ni3P precipitates, which have different crystal structures [6]. However, some doubts remain about this conclusion because only specimens with a
⁎ Corresponding author. E-mail address: [email protected] (H. Miyamoto).
single P content were used for dark-field transmission electron microscopy observations of Ni3P precipitates. Therefore, it is necessary to examine the strengthening mechanism in more detail using specimens with a wider range of precipitate volume fractions. There are currently several proposals for the deformation mechanism in nanocrystalline materials, such as grain boundary sliding and the formation/disappearance of dislocations at grain boundaries. For this reason, the strengthening mechanism in Ni–P composites is an extremely interesting subject of research, because it is likely to provide important insights into the deformation mechanism, and thus greatly contribute to the future development of nanomaterials. In the present study, the strengthening mechanism in Ni–P alloys was examined by forming Ni–P films by electrodeposition, and then producing ultrafine, two-phase Ni–Ni3P alloys by aging under various conditions. 2. Experimental methods In this study, Ni–P alloy specimens were produced by electrodeposition. Table 1 shows the bath composition and electrodeposition conditions used. The P content in the alloys was changed by varying the amount of H3PO3 added to the bath, and H3PO4 was also added to adjust the bath pH to 1.5. The anode and cathode substrates used during electrodeposition were Ni (99.9%) and stainless steel (SUS 304) having a size of 20 × 30 mm, respectively, and they were ultrasonically cleaned in acetone prior to use. The bath was stirred constantly during electrodeposition using a magnetic stirrer, in order to homogenize its composition and temperature, and to withdraw
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Ni(SO3NH2)2–4H2O [g/ℓ] NiCl2–6H2O [g/ℓ] H3BO3 [g/ℓ] H3PO3 [g/ℓ] Plating temperature [K] Current density [mA/cm2] Type of current Plating time [h]
P content [mass%]
Table 1 Bath composition and deposition conditions. 470 15 10 1, 3, 5, 7, 10, 15 323 200 D.C. 2
10
5
0 0 the gas generated by ionization away from the cathode surface to prevent pitting. Deposition was carried out using a pulsed power supply (HCP-301H, Hokuto Denko Corporation). The electrolyte used was a mixed liquid comprising 40% acetic acid, 30% phosphoric acid, 20% nitric acid, and 10% distilled water. As shown in the thermal equilibrium diagram for Ni–P alloy in Fig. 1 [17], the amount of P in solid solution in Ni is greatest (0.32 at.%) at around 900 °C, and P hardly forms a solid solution at room temperature. Therefore, the as-deposited film is a supersaturated solid solution, and aging causes precipitation of Ni3P. In this experiment, the aging treatment involved heating the specimens under a vacuum for 40 min at temperatures of 473 to 773 K. X-ray diffraction (XRD; SmartLab, Rigaku Corporation) analysis was performed in order to determine the crystal structure and the grain size in the specimens. The diffraction conditions were a voltage of 40 kV, a current of 200 mA using CuKα, a scan speed of 4°/min, and an angular range of 30 to 80°. Hardness measurements were carried out using a micro-Vickers hardness tester (HMV-2200, Shimadzu Corporation) under constant load of 980.7 mN. Microstructural observations were performed using transmission electron microscopy (TEM; JEM-2100F, JEOL Ltd.) at an accelerating voltage of 200 kV, and local compositional measurements were carried out using energy-dispersive X-ray spectroscopy (EDS) in the TEM system. The P content in the specimens was determined using inductively coupled plasma optical emission spectrometry (ICP-OES).
3. Experimental results Fig. 2 shows the relationship between the amounts of H3PO3 in the bath and the P concentration after electrodeposition, as determined by
L
1200
31
800 0.32
880
19
(Ni)
Ni2P
400
Ni3P Ni5P2 Ni12P5
600
200
850 825 770
NiP3
870
G
700 450
NiP2
23.5
1000
Ni5P1 Ni P 1.22
Temperature [ ]
1600 1455 1400
(P)
0 0
Ni
10
20
30
40
50
60
70
80
at%P
Fig. 1. Thermal equilibrium phase diagram for Ni–P [17].
90 100
P
5
10
15
20
Fig. 2. Dependence of phosphorous content in as-deposited Ni–P on amount of H3PO3 added.
ICP-OES. The P content is seen to increase almost linearly with the amount of H3PO3 as was reported elsewhere [7]. Fig. 3 shows XRD patterns for as-deposited and aged specimens with different P contents. When the P content is low, the peaks are sharp before and after the heat treatment, showing that the specimens are polycrystalline. However, with increasing P content, peak broadening is observed for the asdeposited specimens, indicating the presence of an amorphous state. The heat treatment causes crystallization to take place, as evidenced by sharpening of the peaks. The temperature at which this occurs tends to increase with P content. For the specimen with 0.8 wt.% P, only peaks associated with Ni were identified. However, for P contents of 5.0, 6.1, 7.7, and 12.5 wt.%, Ni3P peaks were apparent for specimens aged at 673 K. Fig. 4 shows the effect of the P content on the grain size in the Ni parent phase, as determined from the XRD results using the Scherrer formula. For the as-deposited specimens and those heat treated at 473 K, the grain size decreased with increasing P content. However, for the specimen heat treated at 673 K, the grain size was independent of P content. Fig. 5 shows the relationship between the aging temperature and the micro-Vickers hardness. For 0.8%P, the hardness increased monotonically with aging temperature whereas the hardness of higher P contents increased with aging temperature, with the maximum hardness being obtained at 673 K. Figs. 6 to 8 show bright-field, dark-field TEM images, and selected area electron diffraction patterns (SADP) for specimens with a P content of 5.0 mass% after electrodeposition, and aging treatment at 473 K and 673 K as an example. In the diffraction pattern in Fig. 6(b) for the asdeposited specimen, only Debye rings associated with Ni were identified. The dark-field image in Fig. 6(c) was taken using the region of the Debye ring indicated by the circle in the diffraction pattern. From dark-field observations, the average Ni grain size was determined to be 5 nm. For the specimen heat treated at 473 and 573 K, the SADP again contained only Debye rings associated with Ni, and the average grain size determined from the dark-field images was 5 and 9 nm, respectively. For specimens heat treated at 673 and 773 K, diffraction spots associated with Ni and Ni3P were identified, and the average grain size for the two phases was 25 and 23 nm, respectively, for the specimen heat treated at 673 K, and 85 and 74 nm, for the specimen heat treated at 773 K. Fig. 9 shows a compilation of TEM results for specimens with different P contents heat treated at 673 K, which is the temperature that yielded the maximum hardness. For a P content of 0.8 wt.%, no Ni3P was observed, and the average Ni grain size was 35 nm. For a P content of 3.0 wt.% and above, Ni and Ni3P were observed, and the average grain size of both Ni and Ni 3P was about 25 nm, which is comparable to that calculated using the Scherrer formula.
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(200) (111) (200)
(111)
(220)
(a) 0.8 mass%
(220)
(b) 3.0 mass% (220)
(111) (200)
(220)
(c) 5.0 mass%
(111)
(200)
(d) 6.1 mass%
(111) (200)
(220)
(e) 7.7 mass%
(f) 12.5 mass%
Fig. 3. XRD patterns for as-deposited and aged specimens with P contents of (a) 0.8 mass%, (b) 3.0 mass%, (c) 5.0 mass%, (d) 6.1 mass%, and (e) 12.5 mass%.
4. Discussion 4.1. Comparison with rule of mixtures Assuming that all P is precipitated as Ni3P after the aging treatment, the volume fraction of Ni3P will be 6.0 to 95.4 vol.% depending on the P
content. Therefore, when considering the strengthening mechanism for the Ni–Ni3P alloy, the rule of mixtures for Ni and Ni3P cannot be ignored. Fig. 10 shows the values obtained by the experiments and those predicted by the rule of mixtures for specimens aged at a temperature of 673 K. The Ni grain size was taken to be the average of the grain sizes estimated using the Scherrer formula and measured directly from the dark-field
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1200
40
0.8 P (mass%) 3.0 5.0 7.7
30
Vickers hardness [HV]
Grain size [nm]
35 25 20 15 10 5 0 0
2
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6
8
10
12
1000
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P contents [mass%]
600 400 200 0 200
(a) As-deposit
300
400
500
600
700
800
900
Aging temperature [K]
40 35
Grain size [nm]
800
Fig. 5. Dependence of micro-Vickers hardness on aging temperature.
30 25 20 15 10 5 0 0
2
4
6
8
10
12
14
P contents [mass%]
Vickers hardness values for Ni and Ni3P with a grain size of 25 nm were determined to be 530 and 950 HV, respectively. The predicted Vickers hardness for different Ni3P volume fractions is indicated by the dashed line in Fig. 10. It can be seen that the measured values are larger than the predicted values for all compositions. Therefore, it is possible that some additional strengthening mechanism occurs in the Ni– Ni3P alloy.
(b) 473 K 4.2. Grain size distribution in Ni–Ni3P alloy
40
Grain size [nm]
35
Age hardening of Ni–P alloys has often been explained as classic precipitation hardening due to the presence of Ni3P, similar to that which
30 25 20 15
(a)
10
(b)
5 0 0
2
4
6
8
10
12
14
P contents [mass%]
(c) 673 K Fig. 4. Dependence of grain size on P content for (a) as-deposited specimens, (b) specimens heat treated at 473 K, and (c) specimens heat treated at 673 K.
100nm
(c) TEM images. The average grain sizes for specimens with a P content of 0.8, 3.0, 5.0, 6.1, 7.7, and 12.5 wt.% were 35, 24, 25, 26, 26, and 24 nm, respectively. Since the grain size of the specimen with the 0.8 wt.% P content was exceptionally large, the hardness value was excluded when comparing with the rule of mixtures. The Vickers hardness for single phase Ni and Ni3P with a grain size of 25 nm was calculated using the method reported in Ref. 11, assuming that the Hall–Petch law was valid. The Hall–Petch coefficient k is reported to be proportional to the square root of the shear modulus [8]. Therefore, k for Ni3P was determined from the known value for Ni, using the ratio of the shear moduli of the two materials. Using a value of 64 GPa for the shear modulus for Ni3P [18], the relationship kNi3P = 0.92kNi was obtained, and the
100nm Fig. 6. TEM results for as-deposited specimens with a P content of 5.0 mass%: (a) brightfield image, (b) diffraction pattern, and (c) dark-field image of Ni grains obtained from the region indicated by the circle in the diffraction pattern.
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(a)
(b)
5
(a)
(b) A B
100nm
(c)
100nm
(c)
100nm
(d)
100nm
100nm
Fig. 7. TEM results for specimens with a P content of 5.0 mass% heat treated at 473 K: (a) bright-field image, (b) diffraction pattern, and (c) dark-field image of Ni grains obtained from the region indicated by the circle in the diffraction pattern.
Fig. 8. TEM results for specimens with a P content of 5.0 mass% heat treated at 673 K: (a) bright-field image, (b) diffraction pattern, (c) dark-field image of Ni grains obtained from region A in the diffraction pattern, and (d) dark-field image of Ni3P grains obtained from region B in the diffraction pattern.
occurs in materials with much larger grains. However, for this to be the case, the Ni3P grains would need to be considerably smaller than the Ni grains, and to exist within their interior. Otherwise, the Ni3 P grains could not impede dislocation slip. Therefore, the grain size distribution for Ni and Ni3P was determined from TEM dark-field images of specimens that had been heat treated at 673 K, which is the temperature that produced the maximum hardness. Fig. 11 shows the Ni and Ni3P grain size distributions for specimens with P contents of 3.0, 5.0, and 6.1 wt.%. Although a wide distribution is obtained for all specimens, there is no significant difference between the average and median values for Ni and Ni3 P. Therefore, although there may be very small second-phase particles partially dispersed inside the Ni grains, it is considered unlikely that precipitation hardening is the dominant strengthening mechanism. One possibility is that strengthening occurs due to the heterophase interface between the cubic Ni grains and the tetragonal Ni 3 P precipitates [19]. The extra increase of hardness over the rule-of-mixture prediction, which already takes into account the strengthening effects of nanocrystalline structure, must be related to the Ni–Ni3 P interphase boundaries. It is rational to consider that the interface between dissimilar phases (cubic–tetragonal) is more effective than homophase grain boundaries in impeding slip transmission across the interface [19]. The Burgers vectors for dislocations on either side of the interphase interface are different, so the interface may strongly suppress dislocation movement, in contrast to the case for a single-phase grain boundary. The higher density of “geometrically necessary dislocations” at the interface as a result of the inherently greater strain incompatibility between adjacent grains of dissimilar materials can also make the interface harder [19,20]. In Fig. 10, it can be seen that the largest differences between the measured and predicted hardness values are obtained when the volume fractions of Ni and Ni3P are roughly the same. In such a situation, the heterophase interface makes up the largest fraction of the total interface area, and would
therefore be expected to have the largest effect on dislocation movement. Both the strengthening mechanism argued by the present authors and by Chang et al. [11] exclude dislocation pinning by Ni3P precipitation as the primary role. However, there is a clear difference. In Chang's case, the change from non-equilibrium to equilibrium grain boundaries during the annealing is the cause of the strengthening. In as-electrodeposited state, grain boundaries are at non-equilibrium state with extrinsic grain boundary dislocations and trapped extra dislocations. These non-equilibrium grain boundaries can operate as dislocation sources in deformation, and as a result they become transplant for crossing dislocations. Thus, the relaxation at non-equilibrium grain boundaries on annealing makes it more difficult for them to dislocate. The so-called “hardening by annealing” phenomenon has been reported for metals by electrodeposition [21,22] and severe plastic deformation [23–25]. However, it is important to note that the degree of strengthening by annealing in electrodeposited and severely deformed pure metals was no more than 10%, and was much smaller than that in the present study and that reported by Chang et al. [11]. Therefore, the effect of structural change from non-equilibrium to equilibrium state of grain boundaries on the strengthening should be limited in the Ni– P system though it is not negligible. Thus, it is rational to consider that P atoms either as solute or second phase plays a primary role in the strengthening upon the annealing. Although it has been argued that all P is precipitated as Ni3P, it is also possible that some P segregates at grain boundaries [16]. This may also influence the material strength, because grain boundaries are involved in the deformation mechanism for nanocrystalline materials. Fig. 12 shows a TEM–EDS analysis of a specimen with a P content of 5 wt.%, that was aged at a temperature of 673 K. The Ni and P contents were measured along the line profile indicated in red. It can be seen that the P signal is strong within regions with a width of about 30 nm, which corresponds to the size of Ni3 P precipitates.
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P=0.8 mass%
P=3.0 mass%
P=6.1 mass%
P=7.7 mass%
P=12.5 mass%
(a) 100nm
100nm
(b)
A
B
A B
100nm
100nm
100nm
B
A B
B
A
A
(c) 100nm
100nm
100nm
100nm
100nm
100nm
100nm
100nm
100nm
100nm
(d)
Fig. 9. TEM results for specimens heat treated at 673 K: (a) bright-field images, (b) diffraction patterns, (c) dark-field images of Ni grains obtained from region A in the diffraction patterns, and (d) dark-field images of Ni3P grains obtained from region B in the diffraction patterns.
Also, the sizes of the regions with strong P and Ni signals are comparable,. However, there is no evidence for localized P signal at grain boundaries. This indicates that the amount of segregation of P to grain boundaries is small, and its effect on strengthening can be considered low.
P content [mass %] 0
5
10
Vickers hardness [HV]
1200 1000 800 600 400
Ni-Ni3P experimental value Rule of mixture of Ni-Ni3P (d=25µm)
200
5. Summary To clarify the age hardening mechanism in Ni–P alloys, specimens with different P contents formed by electrodeposition were subjected to aging treatments at different temperatures. The hardness was found to increase with aging temperature, and the amount of Ni3P precipitated increased with P content. Regardless of the P content, the maximum hardness was obtained for specimens aged at 673 K, and for these specimens, the Ni and Ni3 P grain size distributions were evaluated by X-ray diffraction and electron microscopy. The results showed that the grain size distribution was similar for Ni and Ni3P. This makes it unlikely that precipitation hardening, in which dislocation movement is hindered by small particles within grains, plays a significant role. In addition, the hardness after aging was higher than that predicted using the rule of mixtures for all Ni3P volume fractions. The largest differences occurred when the volume fractions of the two phases were almost the same. This suggests the possibility that the interface between the Ni and Ni 3P phases plays a large role in the strengthening mechanism.
0 0
20
40
60
80
100
Acknowledgments
Volume fraction of Ni3P [vol%] Fig. 10. Measured Vickers hardness and that predicted by the rule of mixtures.
The authors gratefully acknowledge the financial support of a Grant-in-Aid for Scientific Research on Innovative Areas “Bulk nano
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(a) 0.25 N=100
Frequency
0.2 0.15
Ni Ni3P
0.1 0.05 0
Grain Size [nm]
(b) 0.25 N=100
Frequency
0.2
Fig. 12. TEM image and EDS line profiles for Ni and P in a specimen aged at 673 K. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
0.15 0.1
Ni Ni3P
0.05 0
Grain Size [nm]
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(c)
[8]
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0.2 0.15
Ni Ni3P
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[9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
0.05 0
Grain Size [nm] Fig. 11. Ni and Ni3P grain size distributions in specimens aged at 673 K, with P contents of (a) 3.0 mass%, (b) 5.0 mass%, and (c) 6.1 mass%.
[19] [20] [21] [22] [23] [24]
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