金属硼氢化物基固态储氢体系

顾婷婷; 顾坚; 张喻; 任华

doi:10.7536/PC190829

化学进展 ›› 2020, Vol. 32 ›› Issue (5) :665-686 DOI: 10.7536/PC190829

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金属硼氢化物基固态储氢体系

顾婷婷 ¹^,² ,
顾坚 ¹^,²^,^** ,
张喻 ³ ,
任华 ¹^,²

作者信息 +

Metal Borohydride-Based System for Solid-State Hydrogen Storage

Tingting Gu ¹^,² ,
Jian Gu ¹^,²^,^** ,
Yu Zhang ³ ,
Hua Ren ¹^,²

Author information +

文章历史 +

摘要

氢气储存仍是制约氢经济推行的关键问题,开发一种高效、安全的储氢技术仍面临着巨大挑战。近年来,利用固态氢化物的化学吸附储氢技术由于可靠、结构紧密和高储氢容量的特点,被视为最有潜力的储氢手段之一。在众多固态氢化物储氢材料中,金属硼氢化物由于其极高的重量和体积储氢密度而备受关注。然而,金属硼氢化物热力学稳定,动力学缓慢,导致其吸/放氢温度高、速率慢、可逆性及循环稳定性差。本文从替代、复合、掺杂、纳米结构限域及相应的反应机理等角度总结了金属硼氢化物储氢材料的最新改性研究和应用,并提出了其中存在的问题和相应对策,同时指出了未来的研究方向。

Abstract

Hydrogen storage is the key technological problem for a viable hydrogen economy and so far, finding an efficient and safe method of storing hydrogen remains an indomitable challenge. Chemical sorption via solid-state hydrides, offering reliable, compact and high capacity features, is considered one of the most promising avenues for hydrogen storage. Among the diverse hydrides, metal borohydrides are excellent candidates on account of their high gravimetric and volumetric density. However, these hydrides generally suffer from high temperatures of de/rehydrogenation, slow sorption rate, limited reversiblity and poor cyclability due to the intrinsic thermodynamic stability and/or sluggish kinetics. In this review, we summarize recent researches and applications on the aspect of optimizing performance through substitution, composite, doping, and nanostructure, cognizing the relevant reaction mechanism for the metal borohydride-based system. The challenges and countermeasures are illustrated, and the direction to further enhancing the hydrogen storage properties of the system is also pointed out.

Contents

1 Introduction

2 Metal Borohydrides

2.1 Substitution

2.2 Composite

2.3 Doping

2.4 Nanostructure

3 Conclusion and prospect

Graphical abstract

关键词

储氢材料 / 热力学 / 动力学 / 金属硼氢化物 / 性能调控

Key words

hydrogen storage materials / thermodynamics / kinetics / metal borohydrides / performance adjustment

引用本文

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顾婷婷, 顾坚, 张喻, 任华. 金属硼氢化物基固态储氢体系[J]. 化学进展, 2020, 32(5): 665-686 DOI:10.7536/PC190829

Tingting Gu, Jian Gu, Yu Zhang, Hua Ren. Metal Borohydride-Based System for Solid-State Hydrogen Storage[J]. Progress in Chemistry, 2020, 32(5): 665-686 DOI:10.7536/PC190829

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1 引言

众所周知,化石燃料驱动的汽车尾气是导致当今温室气体排放超标的原因之一^[1]。面对化石燃料殆尽及其燃烧产物导致环境污染的双重压力,各国政府颁布了严苛的法律和政策来维护地球的环境^[2]。

当传统燃料车辆退出历史舞台^[3],为了应对能源结构的突变,寻找一种替代能源显得尤为迫切。目前,多种可再生能源已被开发,比如风能、太阳能、潮汐能等^[4],然而这些能源具有间歇性和地域性,不适用于“移动”应用场景^[5]。锂离子电池^[6]、超级电容器^[7]、燃料电池^[8]是近年来开发的“新能源”技术。不同于锂离子电池和超级电容器储存电能的特性,燃料电池以甲醇、肼、碳氢化合物、氢等为原料,通过原料在电极处的氧化还原反应产生电能^[8]。其中,氢气燃料电池具有最优异的综合性能。因为氢元素是宇宙中最丰富的元素,氢气燃料电池的产物只有水且化学能高达142kJ/g^[9,10],因此,氢气燃料电池的应用被认为是电动车辆发展的终极目标^[11]。

氢气并不是一种能源,而是一种能量载体,其体积密度仅为汽油的1/2700,为了方便使用,需以一种合适的形式存储^[12]。通常氢气采用高压及低温液化的方式存储,但是高压存储的方式安全性差,体积密度低,70 MPa下仅能存储0.03 kg/L氢气^[13];通过低温液化储氢的方式能够获得相对较高的储氢密度(0.07 kg/L),但液化须消耗大量能量,且使用过程中面临着液化氢气蒸发、绝热和热传递等更严峻的考验^{[14, 15]}。因此,氢的高效存储仍然是其大规模应用中的一大难题。

对于车载应用而言,储氢介质应轻质且密实,需具备以下特点:高储氢容量、快速反应动力学、高度可逆性、合理的成本及安全性等^{[16, 17]}。显然,传统的储氢手段无法满足移动化的应用需求。20世纪70年代初,荷兰Philips公司和美国Brookhaven实验室先后发现了LaNi₅、Mg₂Ni等合金具有可逆吸放氢气的能力,其通过化学键将H原子“束缚”在固态储氢材料中,一举解决了传统储氢方式氢气泄漏和安全性差的问题^{[18, 19]},固态储氢材料登上历史舞台。为了促进车载固态储氢材料的研究发展,美国能源部(DOE)根据燃料电池汽车(FCVs)的设计和性能参数,提出了车载储氢材料标准,见表1 ^[16]。

为了满足上述要求,研究者们设计开发了多种固态储氢材料,根据氢的键合方式可分为物理吸附和化学结合材料。对于物理吸附储氢材料,氢借助范德华力以分子形式吸附在储氢材料表面。为了获得大的储氢容量,这些材料需具有大的比表面积、高的微孔体积分数及高的表面吸附能力^{[20, 21]}。研究发现,较碳材料和金属-有机骨架(MOFs),晶体型共价键有机骨架(3D-COFs)有更高的比表面积,因而具备更高的储氢容量^{[22, 23]}。通过元素修饰,如锂、钠、钾和铝等,氢气与MOFs或碳之间的稳定性得到增强^[24],结合能提高至0.1~0.2 eV,增加了物理吸附材料在室温下的储氢容量^[25]。通过结构优化,物理吸附材料的室温储氢性能得到了改善,但由于氢与吸附材料间的作用力仍较弱,使得每单位重量材料的吸氢量仍仅有1 wt%^[26]。物理吸附多孔储氢材料的另一个痛点是其较低的体积储氢密度^[27]。研究发现,虽然物理吸附材料的质量储氢密度在不断提升,但其体积容量并未得到明显改善^[28],而物理吸附储氢系统的体积密度比重量密度对燃料电池汽车的行驶距离影响更大^[29]。

化学固态储氢材料,由于高的储氢密度和稳定的室温存储性,被视为是燃料电池的理想氢源材料。已报道的化学固态储氢材料主要有以下3种:

(1) 金属间化合物,可表示为A _mB _nH _x,其中A和B都是金属,其储氢能力取决于间隙氢原子和金属原子之间的作用力,此类储氢材料性能与其晶体结构关系密切^[30]。迄今为止,研究者已开发如下5种金属间化合物,包括AB₅(CaCu₅型)、AB₂(Laves相)、AB(CsCl型)、A₂B(AlB₂型)和钒基固溶体合金。上述材料较物理吸附材料具有更大的储氢容量^[31]、合适的热力学性能(ΔH=25~50 kJ/mol H₂),在合适的吸放氢温度区间能可逆吸放一定量的氢气。然而,金属间化合物含有重金属及其合金,导致其储氢容量仍小于3 wt%,难以满足车载储氢材料的需求。此外,缓慢的吸放氢动力学性能以及复杂的预活化过程,亦限制了其在移动领域的应用^[32]。

(2) 金属氢化物,代表材料为AlH₃和MgH₂,其质量储氢容量较高,分别为10.1 wt%和7.6 wt%。AlH₃的放氢温度处于130~160 ℃,能够释放约8.3 wt% H₂,接近储氢材料的理想放氢工作区间。然而,AlH₃的放氢是放热反应,其放氢产物Al无法再次氢化,限制了其进一步应用^[33]。MgH₂能够可逆吸放7.6 wt%的氢气,但其反应焓高(约为75 kJ/mol H₂)^[34]、吸放氢动力学缓慢^[35],且易氧化^[36],这亦阻碍了其进一步应用。为了解决MgH₂的热力学和动力学性能缺陷,研究者开展了多种改性研究,如添加催化剂和其他化合物^[37]、纳米化Mg单质^[38]以及利用多孔材料将其纳米限域^[39]等。但非氢相的引入和易氧化的本质极大降低了MgH₂基材料的储氢容量,使其储氢容量不再具备优势^[40]。

(3) 配位氢化物,如铝氢化物、氨基化物和硼氢化物等,虽具有较高的储氢容量,但在研究之初并不是储氢材料的首选,因为其热力学稳定、动力学缓慢且可逆性差。1997年,Bogdanovi 和Schwickardi研究发现,添加少量Ti(OBu ⁿ)₄可使NaAlH₄在温和条件下实现可逆吸放氢^[41],由此引发了将轻金属配位氢化物作为储氢材料的研究热潮。催化改性后,NaAlH₄的热力学性能和动力学性能均得到了改善,但该体系经100次循环后(150 ℃吸氢和160 ℃放氢),仅有3.5 wt% H₂被可逆存储,远低于美国DOE的设定目标^[42]。其他金属铝氢化物,如LiAlH₄和Mg(AlH₄)₂具有更高的储氢容量(分别为10.5 wt%和9.3 wt%),但其放氢过程均为放热反应,放氢产物无法再次吸氢^{[43, 44]}。2002年,Chen等首次报道了Li₃N能够吸收11.4 wt%的H₂,生成LiNH₂和LiH,进而将固态储氢材料的研究拓展至氨基化物^[45]。借鉴于MgH₂的热力学调控方式,研究者们设计了Li-Mg-N-H和Li-Ca-N-H体系以期获得更优的吸放氢性能^[46],但即便是通过成分调变^[47]、催化掺杂^[48]或纳米化^[49],仍无法获得合适的储氢容量,副产物氨气的生成更导致了循环储氢容量的衰减^[50]。2003年,Züttel等发现了LiBH₄具有储氢能力,且具有很高的储氢容量,随后,越来越多的研究者把金属硼氢化物视为最具潜力的固态储氢材料之一^[51],其详细介绍见第2节。

2 金属硼氢化物

金属硼氢化物通式为M(BH₄) _n(n为金属M的价态),其具有极高的储氢容量,是固态储氢材料的研究重点^[51]。此类化合物中氢和硼通过共价键相连形成[BH₄]^-,所带负电由Li、Na、K、Be、Mg、Ca等金属阳离子来补偿^{[52, 53]}。第一性原理计算显示,M(BH₄) _n的电子结构为非金属结构,能隙较大,约为1.8~6.8 eV^[54],M(BH₄) _n的稳定性由

M n +

向[BH₄]^-的电子传输能力决定^[55]。表2给出了几种代表性金属硼氢化物的储氢性能;由于强共价键的存在,LiBH₄、NaBH₄、Mg(BH₄)₂和Ca(BH₄)₂的放氢温度较高;Be(BH₄)₂、Al(BH₄)₃和Mn(BH₄)₂拥有较高的反应活性及储氢容量,但氢气释放的同时伴随大量B₂H₆的产生。NaBH₄亦可通过水解的方式放氢,反应放热,生成NaBO₂,但不再具有可逆吸氢的能力^[66]。高储氢容量、合适的热力学动力学性能和高度可逆性是M(BH₄) _n作为储氢材料的必要条件。因此,本文主要聚焦于通过热解方式且具备可逆吸放氢能力的M(BH₄) _n的最新研究进展,如LiBH₄、NaBH₄、Mg(BH₄)₂和Ca(BH₄)₂等。

金属硼氢化物的分解放氢过程可粗略概括为以下两个反应:

(1)M(BH₄) _n→nMH+nB+3/2nH₂

(2) nMH+nB→nB+nM+1/2nH₂

在不同的反应条件下(不同的氢气背压^[67]、加热速率^[68]、保温时间^[69]等),金属硼氢化物的放氢产物不尽相同^{[69,70,71,72,73]}。最常见的两种产物为最终产物硼和中间产物M(B₁₂H₁₂) _n/ ₂。两者均为二十面体结构,与初始储氢材料的四面体结构具有明显差异^[74]。放氢产物晶体结构的巨大变化阻碍了其在吸氢过程中通过原子重排向初始材料晶体结构的转变,进而限制了该类材料的热力学、动力学及可逆性^[75]。

通过吉布斯自由能的变化(ΔG)可以判断反应是否能够自发进行,其表达式如下:

(3) ΔG=ΔH-TΔS

其中,ΔH为反应焓变,ΔS为反应熵变。在吸放氢过程中,熵变ΔS主要来源于气态氢(氢的熵值为

S H 2 0

=130 J/K mol^[76]),换句话说,反应焓变的差异直接决定了体系的热力学稳定性。M(BH₄) _n放氢过程的焓变表达式如下:

(4) ΔH _desorb=ΔH _form-prod-Δ

H form ⁃ M (B H 4) n

其中,ΔH _form-prod为产物的标准生成焓,Δ

H form ⁃ M (B H 4) n

为理论计算所得M(BH₄) _n生成焓。对于可逆吸氢反应,放氢须为吸热反应,即ΔH _desorb>0,而产物生成焓(ΔH _form-prod)和M(BH₄) _n生成焓(Δ

H form ⁃ M (B H 4) n

)均小于零。因此,通过稳定化放氢产物(ΔH _form-prod更负)或对金属硼氢化物去稳(Δ

H form ⁃ M (B H 4) n

由负及零),体系即可获得更小的放氢反应焓,实现体系在温和条件下的吸放氢,其具体调控原理见图1。

基于金属硼氢化物的基本信息,本文总结了金属硼氢化物改性研究的最新进展:(1)热力学调控,包括替代和复合;(2)动力学调控,通过催化掺杂或原位反应生成的活性物质,提供足够活性反应点来提高反应动力学;(3)纳米化,通过减少颗粒尺寸、优化材料形貌和纳米限域的方式,提高体系热力学和动力学性能。

2.1 替代

如上所述,从M ⁿ ⁺向[BH₄]^-的电子传输能力是决定M(BH₄) _n热力学稳定性的关键因素。进一步通过第一性原理计算发现,M(BH₄) _n的生成焓Δ

H form ⁃ M (B H 4) n

和

M n +

的泡利电负性χ_p密切相关,χ_p的差异会直接影响M(BH₄) _n的热力学稳定性^[77]。因此,金属硼氢化物的放氢焓变ΔH _desorb可由χ_p预测,表达式如下:

(5) ΔH _desorb≈248.7χ_p-M ⁿ ⁺-390.8

换言之,放氢温度T _d和χ_p之间也有相似的线性关系^[78],通过调控金属阳离子的电负性即可获得可期的体系工作温度。然而,单一金属仅可提供有限而固定的χ_p,无法灵活调节硼氢化物的热力学稳定性。对此,研究者在单金属硼氢化物的基础上,设计合成了双阳离子硼氢化物MM'(BH₄) _n,其通过调整金属复合物的χ_p来实现体系热力学的调控。

2.1.1 双阳离子硼氢化物

通常MM'(BH₄) _n由以下两种方法合成:(1)硼氢化物和金属卤化物反应^[79],如式(6);(2)两种硼氢化物球磨^[80],如式(7):

(6)MCl _n+mM'BH₄→MM _m-n(BH₄) _m+nLiCl

(7)LiBH₄+xM(BH₄) _n→LiM _x(BH₄

) 1 + n

方法1将产生40 wt%~60 wt%的LiCl,显著降低了体系的储氢容量。引入合适溶剂能够实现硼氢化物和LiCl的分离,但去除效果十分有限,反而又会引入含醚基的副产物^[81],进一步降低了体系的储氢容量。方法2所制双阳离子金属硼氢化物的纯度亦低于90%,且对于部分合成物的成分和结构仍存争议^[82,83]。表3列举了代表性双阳离子硼氢化物的合成方法及其纯度。

最近,研究者提出了一种新的室温湿化学法制备混合金属硼氢化物的方法,其利用溶剂调控体系的离子交换反应,制备出了高纯度的混合金属硼氢化物,具体制备过程如图式1^[94,95]。步骤C,釜中发生机械诱导化学反应,如公式(8)和(9):

(8)M[An]+[Cat][Zn₂(BH₄)₅]→M[Zn₂(BH₄)₅]↓+[Cat][An](M=Li)

(9)M[An]+[Cat][Zn(BH₄)₃]→M[Zn(BH₄)₃]↓+[Cat][An](M=Na, K)

通过上述反应,混合阳离子硼氢化物以沉淀的形式从[Cat][An]溶剂中析出,极易分离,所制Li[Zn₂(BH₄)₅]和Na[Zn(BH₄)₃]的纯度分别高达99.8 wt%和92.4 wt%。

此外,通过在氨气气氛下的反应,不同金属硼氢化物间能够直接反应形成氨稳定的双阳离子硼氢化物,即MM'(BH₄) _n·mN

H 3

^[96]。该方法合成过程简单,但MM'(BH₄) _n·mNH₃无可逆性,首次放氢也伴随着NH₃释放。

目前,双阳离子硼氢化物的研究仍然集中在可控合成合适热力学稳定性产物及其结构表征方面^{[97,98,99,100,101]},储氢性能的系统研究较少且暂无突破性进展。

2.1.2 阴离子取代

除阳离子取代,阴离子取代亦可用来调节硼氢化物的性能,其通过反应物球磨实现[BH₄]^-或H被相同电荷和近似尺寸的阴离子局部取代,获得相应固溶体。研究表明,配位阴离子[BH₄]^-整体可被重卤元素取代,如Cl^-、Br^-或者I^-,而F^-仅可取代[BH₄]^-中的H^-[102]。

热分析显示,阴离子替代所得六方相Li(BH₄

) 1 - x

Cl _x反而具有更高的稳定性,其放氢温度较纯LiBH₄略有升高。同样,LiBr^[103]或LiI^[104]的加入也会使得LiBH₄变得更稳定。理论计算表明^[105],引入电负性更强的F^-能够显著降低原有体系的热力学稳定性。例如3LiBH₄-LiBF₄和9LiBH₄-LiBF₄体系,其分解放氢温度低至80 ℃,较单一组分降低了3倍以上,但其放氢时会同时释放大量硼烷^[106]。

基于Li(BH₄

) 1 - x

Cl _x,研究者将阴离子替代拓展至NaBH₄、Ca(BH₄)₂和Mg(BH₄)₂体系。通过NaBH₄和NaCl球磨形成了稳定的Na(BH₄

) 1 - x

Cl _x固溶体,其稳定性随氯化物取代度的增加而提高^[107]。CaX₂(X=F, Cl)的引入使得Ca(BH₄)₂的放氢温度降低了约10 ℃,这主要源于放氢过程中形成了CaH₂-CaF₂和CaHCl固溶体。值得注意的是,Ca(BH₄)₂阴离子替代体系放氢温度的降低归因于放氢途径的优化而非整个系统的热力学改善^[108]。同样,α-Mg(BH₄)₂-MgX₂(X=Cl, B)体系的放氢温度亦有所下降^[109]。然而,CaI的引入会导致α-Ca(BH₄)₂放氢温度的升高,具体原因尚未明确^[110,111]。

2.1.3 离子替代产物的新应用

综上所述,离子取代对于金属硼氢化物的放氢温度并无明显影响,而重金属或金属卤化物的引入显著降低了金属硼氢化物的储氢容量^[112],因此,离子替代并不是金属硼氢化物储氢性能改善的首选方式。

2007年,研究者发现LiBH₄高温相h-LiBH₄为快离子导体,115 ℃时,Li⁺电导率为10^-3 S/cm^[113],170 ℃以上,Li⁺电导率可达到10^-2 S/cm^[114]。通过第一性原理计算发现,o-LiBH₄(低温相)的Li空位或间隙的迁移能垒与h-LiBH₄(高温相)相当,均为0.1~0.3 eV。因此,两者离子电导率的差异被归因为缺陷形成能的高低^[115]。此外,NaB₁₀H₁₀和Na₂B₁₂H₁₂的高温相(相转变温度分别为87 ℃和207 ℃)具有较高的离子电导率,分别达到~0.01和0.1 S/cm^{[116, 117]}。碳局部替代硼原子,NaCB₁₁H₁₂的电导率在较低温度(107 ℃)下即可提升至0.1 S/cm^[118]。基于上述材料电导率的可调性,研究者将金属硼氢化物的应用拓展至固态电池电解质领域^{[119, 120]}。

在室温下如何保持高温相的稳定对于电导率的提升至关重要。研究发现,通过引入LiCl^[121]或LiBr^[122],LiBH₄的高温相能够在低温下保持稳定。Li₄(BH₄)₃I在300 K时具有2×10^-5 S/cm的锂离子快速电导率^{[114, 123, 124]},这是因为I^-的引入使[BH₄]^-间距离增加并诱发晶格非谐性,进而提升了体系的室温电导率。

除了提升高温相h-LiBH₄在低温条件下的稳定性,研究者从降低低温相向高温相转变温度的角度也开展了变相研究。结果发现,NaBH₄的引入,将与LiBH₄形成共晶化合物,显著降低LiBH₄的相变温度。

相变温度的降低,导致体系的活化能从1.04 eV降至0.79 eV,使xLiBH₄-NaBH₄在低温下(<373 K)的Li⁺电导率较LiBH₄提高了10~100倍^[125]。

最新研究发现,LiBH₄在高压条件下(GPa级)会转变为岩盐型(rock-salt-type)LiBH₄相,与h-LiBH₄不同的是,岩盐型LiBH₄相的Li⁺电导率可通过元素掺杂的方式得到大幅提升。未掺杂和5 mol% Ca掺杂的LiBH₄在4.0 GPa和220 ℃下的离子电导率分别为2.2 × 10^-4和 1.4 × 10^-2 S/cm。Ca掺杂后,Li⁺电导率的提高源于生成了双阳离子LiBH₄-Ca(BH₄)₂固溶体,然而,高压撤销后,电导率的提高也将消失^[126]。

2.2 复合体系

2005年,Vajo等基于2LiBH₄-MgH₂体系首次提出了反应氢化物复合(Reactive Hydride Composites, RHCs)的概念,第二相MgH₂的引入,使LiBH₄放氢产物能够在230~250 ℃和100 bar氢气压力下可逆吸收8 wt%~10 wt% H₂,放氢反应如式(10)所示^[127]:

(10)2LiBH₄+MgH₂→MgB₂+2LiH+4H₂

由式(10)和图2可知,MgH₂的引入将使体系生成更稳定的MgB₂,而非B,LiBH₄反应焓变由72 kJ/mol H₂降至46 kJ/mol H₂。热力学性能的改善,使2LiBH₄-MgH₂体系的吸放氢性能较纯LiBH₄得到了极大提升。基于这一发现,更多的第二相被添加至硼氢化物体系以期控制反应焓变,具体可分为以下几类。

2.2.1 元素、金属氢化物或金属氟化物

通过调控非晶B的引入量,LiBH₄和非晶B形成的复合物能够获得合适放氢速率和分解温度。然而,非晶相B的引入对Ca(BH₄)₂和NaBH₄的分解温度并无明显作用,反而会降低其放氢速率^[128]。Al的添加,能够显著提高LiBH₄原始体系的吸放氢动力学,但体系去稳效果无法维持,首次吸氢后相关效果即消失,且伴随有明显的储氢容量损失^[129,130]。Mg的引入,显著改善了LiBH₄和Ca(BH₄)₂的可逆性,但其吸放氢温度仍较高,需进一步降低^[131,132]。

对于复合体系,金属硼氢化物与金属氢化物(MH _x)的复合吸引了研究者的广泛关注,其一方面实现了对体系热力学性能的调控,使体系具备更优异的可逆性,且含氢相的添加还维持了体系的整体储氢容量。值得注意的是,该复合体系的反应产物和反应途径与反应参数密切相关,例如气氛背压^{[132,133,134,135,136,137]}、反应温度^[138,139]或催化剂等^{[140,141,142]}。

在真空或惰性气氛中,2LiBH₄-MgH₂体系组元分别单独分解,且放氢过程有明显的孕育期,期间伴随中间产物Li₂B₁₂H₁₂的生成,最终生成B,反应过程如式(11)^[133]:

(11)2LiBH₄+MgH₂→2LiBH₄+Mg+H₂→1/6Li₂B₁₂H₁₂+5/3LiH+Mg+19/6H₂→2LiH+Mg+2B+4H₂

而在1 MPa的H₂压力下,孕育期明显缩短,Li₂B₁₂H₁₂的生成被抑制,反应直接生成了更利于可逆吸氢的产物MgB₂,反应如式(12)所示^[134]:

(12)2LiBH₄+MgH₂→2LiH+MgB₂+4H₂

无论有无氢背压,6LiBH₄-CaH₂复合体系均表现出优异的可逆储氢性能,反应如公式(13)所示^[143]:

(13)6LiBH₄+CaH₂↔CaB₆+6LiH+10H₂

由于MgB₂为六方结构,而CaB₆为立方结构,放氢产物晶体结构的差异可能导致了2LiBH₄-MgH₂和6LiBH₄-CaH₂复合体系放氢过程的差异^[144]。

Ca(BH₄)₂-MgH₂体系的放氢过程可能经历如下反应:

(14)Ca(BH₄)₂+MgH₂→CaH₂+2B+Mg+4H₂

(15)Ca(BH₄)₂+MgH₂→CaH₂+MgB₂+4H₂

(16)Ca(BH₄)₂+MgH₂→2/3CaH₂+1/3CaB₆+Mg+13/3H₂

由于MgB₂较B具有更高的生成焓,因此反应(15)较反应(14)在热力学上更易发生^[132]。反应(15)和(16)则存在竞争关系,哪个反应占主导取决于实验条件。理论研究显示,在350 ℃和1 bar H₂压力下,更易形成MgB₂,而在350 ℃和动态真空条件下则更易生成Ca

B 6

^[132]。此外,Ca(BH₄)₂-MgH₂体系无须借助催化剂即可获得较好的局部可逆性,这是因为原位生成的Mg可作为体系放氢产物CaB₆的异质形核核心^[144]。进一步地,实验证实Ca(BH₄)₂-MgH₂复合体系在循环过程中会逐渐积累难以可逆的CaB₁₂H₁₂,这也解释了上述体系仅具有局部有限可逆性^[69]。

图2为吸氢van't Hoff曲线,研究发现,温度低于413 ℃时,吸氢反应按式(17)进行,而不是经历常规的式(18)过程:

(17)MgB₂+2LiH+3H₂→2LiBH₄+Mg

(18)MgB₂+2LiH+4H₂→2LiBH₄+MgH₂

温度高于413 ℃时,体系有两个压力平台,分别代表着LiBH₄和MgH₂的生成。除此之外,图2给出了2LiBH₄-MgH₂体系在不同条件下的反应区域,很明显,区域3具有单一反应平台和高的储氢容量,为最适宜的反应区域^[138]。

最近,研究者系统研究了真空条件下,Ca(BH₄)₂在350 ℃和400 ℃保温后的反应产物。研究发现,在350 ℃和动态真空下,Ca(BH₄)₂将分解为CaH₂和CaB₆纳米晶;而在较高温度下(400 ℃),Ca(BH₄)₂分解将首先生成元素B,但随着保温时间的延长,其可转化为Ca

B 6

^[139]。

最新研究显示,除了反应温度和压力,添加剂尺寸的细化亦会极大影响LiBH₄-MgH₂体系的反应途径。对于引入纳米MgH₂的LiBH₄-MgH₂复合体系,其反应过程将历经3个过程(A,B和C),而添加常规MgH₂的LiBH₄-MgH₂体系仅为两步反应,见图3。反应步骤B将生成MgB₂,并显著抑制Li₂B₁₂H₁₂的生成,两者均将显著提高体系的可逆性^[140]。类似地,用纳米尺度的LiBH₄取代LiBH₄后,LiBH₄-MgH₂复合体系的反应路径发生改变,LiBH₄在150 ℃时先于MgH₂分解,释放出大量氢气^[141]。除此之外,引入纳米化的Li _xTiO₂后,将大大缩短LiBH₄-MgH₂体系的放氢孕育期,抑制可逆性差的中间相Li₂B₁₂H₁₂的生成,直接生成了利于再吸氢的Mg

B 2

^[142]。

近来,有研究者在非金属载体石墨烯(GR)和氮化碳(g-C₃N₄)上制备了过渡金属Ni纳米颗粒,引入至2LiBH₄-MgH₂中,经Ni@g-C₃N₄掺杂样品的反应孕育期几乎消失,初始放氢温度降低至105 ℃,储氢容量达9.5 wt%,且LiBH₄放氢反应的活化能降低至126 ± 7 kJ/mol。放氢过程中原位生成了MgNi₃B₂,可作为MgB₂形核中心,促进体系放氢反应。除此之外,g-C₃N₄中的微量H⁺和LiBH₄及MgH₂中的H^-发生反应促进放氢,g-C₃N₄还能够使LiBH₄和MgH₂失稳生成Li₃N和Mg₃N₂,这两种中间产物又易于氢化,增强了体系的可逆性。这种复合催化体系对于低放氢温度和高储氢容量硼氢化物基储氢材料研究具有重要意义^[145]。

除了单金属氢化物,Vajo等提出了引入配位金属氢化物阴离子的复合体系,例如LiBH₄-Mg₂NiH₄的放氢起始温度仅为250 ℃,其具有较低的焓变和熵变(ΔH=(15.4 ± 2) kJ/mol H₂,ΔS=(62.2 ±3) J/K mol H₂)。配位[NiH₄]^4-的改性效果可能源于金属Ni核^[146]。其他过渡金属氢化物阴离子,如[FeH₆]^4-、[CoH₅]^4-和 [MnH₆]^5-也具有类似的改善效果^{[147, 148]}。上述第二相的引入明显降低了LiBH₄的放氢温度,增强了其放氢动力学性能,但低含氢量金属氢化物阴离子的引入降低了体系的整体储氢容量。

为了进一步改善热力学和动力学性能,研究者采用原位生成的M或MH _n,如Al、MgH₂和LaH₃等,对金属硼氢化物进行复合改性,以期从尺寸效应和多组分角度获得协同优化效应。Zhou等通过氢压下机械化学反应的方法,成功制备了LiBH₄-xLa₂Mg₁₇活性复合体系。该体系在氢压下球磨过程中原位生成了MgH₂和LaH₃,两者将协同改性LiBH₄的储氢性能。较球磨外添MgH₂和LaH₃,原位体系的放氢结束温度降低了超过100 ℃。此外,原位体系的放氢产物能够吸收约6.8 wt%的氢气,已接近理论吸氢量(7.6 wt%)^[149]。

LiBH₄中引入Mg(AlH₄)₂后,Mg(AlH₄)₂将首先自分解生成MgH₂和Al,由于原位生成的MgH₂和Al尺寸小,分散度高,且两者会激发协同作用。较单独添加MgH₂或Al的体系,LiBH₄-Mg(AlH₄)₂具有更低的放氢温度,更快的吸放氢速率^[150]。

F^-和H^-结构相似,研究发现,F^-的引入亦会对金属硼氢化物体系的热力学调控发挥积极作用。MgF₂或CaF₂的添加较其对应的金属氢化物(MgH₂或CaH₂)对LiBH₄的放氢性能改善和反应机理并无明显区别,体系可逆性和吸放氢动力学性能较纯LiBH₄均得到了一定改善,但放氢温度并未明显降低^[151]。3LiBH₄-TiF₃复合体系则在70~90 ℃时即可释放出超过5 wt% H₂,但该体系放氢反应放热,反应过程中会生成中间物Ti(BH₄)₃,放氢过程伴随少量B₂H₆释放且产物难以再次吸氢^[152]。最新研究显示,LiBH₄和氟化石墨烯组成的复合体系拥有快速的放氢能力和优异的可逆性,其归因于反应过程中原位生成了LiH _x

F 1 - x

相^[153]。

类似地,对于Ca(BH₄)₂,添加一定量MgF₂后,体系在反应过程中亦会形成中间相Ca

F 2 - x

H _x。Ca(BH₄)₂-MgF₂体系在3次吸放氢循环后,仍保持5.1 wt%~5.8 wt%储氢容量,并仍具有出色的动力学性能^[154]。进一步研究发现,Ca(BH₄)₂的可逆性与氟化物种类有关,Ca(BH₄)₂中引入TiF₄或NbF₅,体系因反应生成CaB₆而具备局部可逆性;而添加了VF₄、TiF₃或VF₃的复合体系只能生成硼元素,放氢产物不具备优良的再吸氢性能^[155]。

为了进一步阐明F^-改性复合体系吸放氢反应机理,研究者从放氢产物出发,系统研究了其吸氢过程,如LiF-Mg

B 2

^[156~159]和CaH₂-MgB₂-Ca

F 2

^[160,161]。基于SR-PXD 结果,MgF₂和LiB

H 4 - x

F _x在氢化过程中就已生成,而Li

H 1 - x

F _x在随后的放氢过程中才生成^[156]。从热力学角度可知,MgF₂由于比MgH₂稳定,其形成促进了体系吸氢过程,而LiB

H 4 - x

F _x活性较LiBH₄高,因而改善了体系的放氢过程^[156]。类似地,循环过程中Ca

H 2 - x

F _x固溶体的生成致使F^-引入的Ca(BH₄)₂体系具有更优异的热力学性能^[160]。

2.2.2 硼氢化物

不同硼氢化物间的复合亦是设计新型储氢体系的有效方法。Lee等首次报道了xLiBH₄-(1-x)Ca(BH₄)₂(0<x<1)复合物体系的可逆性。通过成分调变,NbF₅催化的0.4LiBH₄-0.6Ca(BH₄)₂体系在400 ℃和90 bar氢压条件下,能够可逆存储40%的氢气,明显优于单组分体系^[83]。金属硼氢化物的熔融也会提高反应活性,改善其放氢动力学。经证实,碱金属和碱土金属硼氢化物在高于熔点时分解放氢,其熔点一般高于250 ℃^[162],含LiBH₄的共晶熔融复合体系的放氢温度往往取决于第二组分而不是熔点。在上世纪70年代,实验证明,通过添加NaBH₄形成熔点约为220 ℃的共晶熔融复合物,可使LiBH₄更易分解放氢^[163]。随后,研究开发了多种共晶熔融复合物,见表4。由表可知,共晶体系中,o-LiBH₄到h-LiBH₄的相转变温度亦有明显降低。由于共晶体系分解产物的相互作用,其可逆性得到明显改善^{[170, 171]},但体系的整体性能仍需进一步提高,特别是对于操作温度和吸放氢速率。目前,金属硼氢化物共晶体系的研究仍主要集中在相图研究方面。

值得注意的是,NaBH₄-Ca(BH₄)₂、Ca(BH₄)₂-Mg(BH₄)₂复合体系并没有共晶熔融过程,但仍具有优于单一组分的吸放氢性能^[167,172]。对于不同金属硼氢化物复合体系,通过吸放氢条件的调控,样品熔融过程得到控制,可获得高容量储氢材料。然而,上述系统仍然面临着操作温度高,放氢速率缓慢和循环性能不佳等问题^[173]。

Gu等最新研究表明,三元复合体系Ca(BH₄)₂-2LiBH₄-2MgH₂具有优异的循环性能(10次循环后保持97%的可逆性)。由图4可知,Ca(BH₄)₂-2LiBH₄-2MgH₂体系的放氢结束温度较任意二元体系均低,而储氢容量明显提高。该体系在放氢过程中将原位形成活性双阳离子氢化物(CaMgH_3.72),其显著促进了体系后续的吸氢过程,使得吸氢产物的起始放氢温度降低了约90 ℃^[174]。进一步地,通过氢压下球磨LaMg₃合金,原位引入MgH₂和LaH₃,LiBH₄-Ca(BH₄)₂-xLaMg₃体系的储氢热力学和动力学性能较Ca(BH₄)₂-2LiBH₄-2-MgH₂体系得到了显著提升^[175]。因此,开发多组分储氢体系有望获得综合性能优异的储氢性能。

2.2.3 金属酰胺或氨基络合物

氨基化合物,如LiNH₂,Mg(NH₂)₂和Ca(NH₂)₂中的H⁺,与金属氢化物,如LiH,MgH₂和CaH₂中的H^-相互吸引,驱动了两者复合体系的分解放氢^[46]。考虑到[BH₄]^-中H亦为负电荷,因此,[BH₄]和[NH₂]基团之间亦存在正H和负H作用。Pinkerton等首次报道了2LiNH₂-LiBH₄复合体系,放氢过程形成四元氢化物Li₃BN₂H₈,具有11.9 wt%的理论储氢容量^[176],其在250 ℃下即可释放出超过10 wt% H₂,同时也伴随着少量氨气的释放,反应如式(19):

(19)LiBH₄+2LiNH₂→Li₃BN₂H₈→LiBH₂+4H₂

通过组分调控^{[177,178,179]}、催化掺杂^{[180,181,182,183,184,185,186]}等方法,可进一步改善Li-B-N-H体系的放氢性能并抑制NH₃的产生。Tang等首次报道了钴基催化剂对Li-B-N-H体系放氢动力学的影响,原位形成的活性金属Co和Co₂B改善了体系的储氢性能^[180]。随后,研究者在Li-B-N-H体系中引入了不同Co基化合物,如金属Co^[181]、CoO^[182]、Co₃

O 4

^[183]、Co(OH

) 2

^[184]、Co@C^[185]等,其均能一定程度上改善体系的整体储氢性能。在单纯金属催化的基础上,研究者将金属Ni前驱体与MOF混合,经水热反应,原位生成的Ni纳米颗粒负载于多孔碳上^[187]。较LiBH₄-2LiNH₂体系,LiBH₄-2LiNH₂-10 wt% Ni@C样品的放氢温度降低了110 ℃,放氢动力学显著改善^[187]。

受2LiNH₂+LiBH₄体系启发,研究者开发了多种硼氢化物和氨基化物复合体系,如NaBH₄-NaN

H 2

^[188]、Mg(BH₄)₂-Mg(NH₂

) 2

^[189]、Mg(BH₄)₂-NaN

H 2

^[190]、Ca(BH₄)₂-NaN

H 2

^[190]、Ca(BH₄)₂-LiN

H 2

^[191]、Ca(BH₄)₂-Mg(NH₂

) 2

^[192]、Ca(BH₄)₂-Ca(NH₂

) 2

^[192]、Ca(BH₄)₂-CO(NH₂

) 2

^[193]等。其中,Ca(BH₄)₂-4LiNH₂在250~350 ℃下可释放出超过9 wt%的氢气,比单组分Ca(BH₄)₂显示出更优的放氢性能^[192]。其他Ca(BH₄)₂-氨基化物复合体系的热力学和动力学也得到了类似的增强,但与Ca(BH₄)₂-4LiNH₂相比,其储氢容量相对较低。

除了金属氨基化合物,氨(NH₃)也可以作为H⁺的来源。通过NH₃中H⁺和硼氢化物中H^-的相互作用,氨络合金属硼氢化物M(BH₄) _n·n'NH₃(AMBs)可释放超过10 wt% H₂。然而,由于NH₃络合于M(BH₄) _n,氨络合金属硼氢化物在放氢过程中,NH₃可能会直接释放,而非发生NH₃与[BH₄]^-的相互作用,见式(20)^[193]。

(20)M(BH₄) _n·n'NH₃→M(BH₄)_n·n″NH₃+(n'-n″)NH₃→M(BH₄) _n+n'NH₃

据报道,在温度低于180 ℃的开放空间下,Li(NH₃) _nBH₄会优先释放氨气^[193]。同样,其他AMBs随着温度升高,在开放体系中亦会释放NH₃。因此,在NH₃释放前,氨络合体系必须使N:→M ⁿ ⁺键具有足够的热力学稳定性,以确保BH基团和 NH基团首先发生相互作用,进而释放出氢气。

为了获得高纯H₂,研究者在氨络合金属硼氢化物体系作了大量研究^[194]。Li等通过超声辅助湿化学法成功制备了直径在20~40 nm的Mg(BH₄)₂·6NH₃纳米颗粒,见图5。纳米尺度的Mg(BH₄)₂·6NH₃在未添加任何催化剂的情况下,H₂仍为主要放氢产物^[195]。进一步机理研究表明,纳米Mg(BH₄)₂·6NH₃放氢生成无定形BN和Mg-B-N,反应如式(21):

(21)3(Mg(BH₄)₂·6NH₃)→(6-x)BN+Mg₃B _x

N x + 2

+10NH₃+24H₂

而大颗粒Mg(BH₄)₂·6NH₃放氢产物为Mg和BN^[195]。由于Mg(BH₄)₂·6NH₃体相中的H₂生成能为0.79 eV,高于NH₃分子的生成能(0.67 eV),因此,即使储氢材料的H₂的扩散能垒仅为0.3 eV,其亦难以从Mg(BH₄)₂·6NH₃体相中释放出来。相反,纳米Mg(BH₄)₂·6NH₃的H₂生成能和反应能垒均低于NH₃,使纳米Mg(BH₄)₂·6NH₃的放氢温度降低,且所释放气体的H₂纯度高^[195]。

除了降低颗粒尺寸,引入钴基催化剂亦能显著降低氨络合金属硼氢化物操作温度,同时抑制NH₃的生成^{[193, 196~198]}。此外,基于辅助阴离子/阳离子理论,将AMBs进一步与硼氢化物^{[199,200,201]}、铝氢化物^[202,203]、金属氢化物^[204,205]、金属^[206,207]和氟取代物^[208]等复合,都可以一定程度改善体系的储氢性能,且封闭空间更有利于H₂的生成^[209]。然而,由于上述体系放氢最终生成B-N化合物,无法再次吸氢。尽管通过进一步复合,使氨络合金属硼氢化物基体系获得了一定再吸氢能力,但吸收的H₂均归因于第二相的添加,如Mg

H 2

^[205]。

2.3 掺杂

通过催化掺杂可有效降低金属硼氢化物的活化能垒,显著改善体系的动力学性能。通常,掺杂的催化剂可分为过渡金属和其化合物两类,包括过渡金属氢化物、氧化物、卤化物等。过渡金属卤化物的催化效果要优于其他化合物,原因可归结为三个方面:

(1)固溶体Metal-H-Cl或Metal-H-F的形成。通过含Cl^-和F^-的催化剂掺杂,反应生成新的固溶体,弱化了B—H键,进而提升了体系吸氢性能^[108]。添加TiCl₃后, Ca(BH₄)₂放氢生成Ca-H-Cl固溶体,体系吸氢条件从700 bar氢压和440 ℃降至90 bar 氢压和350 ℃,且掺杂后的样品在相对温和的吸氢条件下能够可逆吸收约3.8 wt%氢气,获得了50%的可逆性^[210]。NbF₅的添加,使Ca(BH₄)₂放氢生成Ca-

H 2 - x

-F _x固溶体,获得了优异的动力学性能和可逆性^[211]。进一步研究发现,F^-对B—H键的弱化作用较其他卤素更强,这也解释了添加含F催化剂储氢性能改善效果优于其他金属卤化物^[212]。理论计算表明,Ti和Nb基催化剂的添加,金属均倾向于占据间隙位点,进而破坏初始的B—H键,促进体系放氢。通过电子局域函数分析,Ti或Nb基催化剂掺杂后,[BH₄]^-基团发生重组,但Ti的催化功能弱于Nb,这与实验结果一致^[213]。还有研究者通过添加氟化石墨烯的方法降低了LiBH₄的放氢温度,提高了其循环稳定性,其亦与F^-替代效应有关^[214]。

(2)过渡金属硼化物的生成。研究者利用X射线吸收光谱(XANES)发现Nb和Ti基催化Ca(BH₄)₂体系在放氢过程中会原位生成TiB₂和NbB₂纳米颗粒。由于晶面间距的临界错配度小于6%,{111}CaB₆/{1011}NbB₂、{111}CaB₆/{1010}NbB₂、以及{111}CaB₆/{1011}TiB₂平面极易配对。因此,原位生成的纳米金属硼化物将作为CaB₆的异质形核核心,促进CaB₆的异相成核并细化微观结构,从而促进Ca(BH₄)₂的吸放氢性能^[215]。

基于金属硼化物的改性效果,研究者直接合成了非晶TiB₂和NbB₂作为催化剂加入到2LiBH₄-MgH₂复合体系中,纳米TiB₂掺杂的2LiBH₄-MgH₂体系在400 ℃保温6分钟即可释放约9 wt% H₂,明显优于未掺杂体系(300 min释放3.9 wt%)。机理研究表明,金属硼化物纳米颗粒首先催化MgH₂的放氢,随后催化LiBH₄分解,并可作为放氢产物MgB₂的成核剂,因此极大促进了2LiBH₄-MgH₂复合体系的放氢过程^[216]。

(3) 活性双阳离子硼氢化物的生成。大量研究表明,过渡金属卤化物的添加对体系储氢性能的改善源于其与金属硼氢化物局部反应生成了相应的硼氢化物。如Zn、Ti和Nb基卤化物添至LiBH₄后,在球磨或放氢过程中会局部生成Zn(BH₄)₂、Ti(BH₄)₃和Nb(BH₄)₅。由于反应生成的过渡金属硼氢化物具有更高的活性,其与LiBH₄形成了混合双阳离子硼氢化物,促进了体系的放氢,但高活性的过渡金属硼氢化物在反应过程中极易释放B₂

H 6

^[217~219]。

除添加上述金属基催化剂,研究者在LiBH₄中引入还原的氧化石墨烯(rGo),使LiBH₄的放氢起始温度得到了进一步降低,这可能是由于反应物会与层状rGo片发生相互作用,改变了原有的分解途径,但其具体的催化机理仍须进一步研究^[220]。除作为基体储氢材料,硼氢化物本身也可以作为其他储氢体系的添加物。LiNH₂可与少量LiBH₄生成一种离子液体相,其促进LiNH₂从固态向熔融态转变,弱化了N—H键,提升了LiNH₂的吸放氢性能^[221]。Hu等发现向Mg(NH₂)-2LiH加入0.1 mol LiBH₄,体系的吸放氢速率较原始Mg(NH₂)-2LiH体系提升了3倍。机理研究显示,反应过程将原位生成Li₃BN₂H₈或(LiNH₂) _x-(LiBH₄)_(1- _x ₎,削弱了N—H键,改善了体系动力学性能^[222]。Li等研究显示,Mg(NH₂)₂-2LiH-0.1Ca(BH₄)₂复合体系具有优异的储氢性能,由于反应过程原位生成了CaH₂和LiBH₄,两者协同作用,使体系能够可逆吸收大于4.5 wt%的H₂,放氢起始温度仅为90 ℃,吸氢温度只有60 ℃^[223]。添加0.1 mol Mg(BH₄)₂的Mg(NH₂)₂-2LiH体系亦获得了类似的改善效果^{[224, 225]}。

2.4 纳米化

就储氢的化学机理而言,调控材料颗粒尺寸和结构亦是一种有效地改善硼氢化物储氢性能的方法^{[226,227,228]}。通过细化颗粒尺寸至纳米级别,或将其制成具有纳米结构的材料,体系将获得新的改性优势,如缩短离子扩散距离,抑制相分离,减小反应界面厚度等,其均将进一步提升体系的动力学性能和可逆性^[229]。众所周知,纳米材料具有高比表面积,因此,表面自由能(W _surf)对体系总能量的贡献无法忽略,可将van't Hoff方程的适用范围拓展至纳米材料,见式(22)和(23)^[230]:

(22)

ln P eq = 1 RT (Δ H 0 + Δ W surf) - Δ S 0 R

(23)

Δ W surf = γ product A product - γ reactant A reactant

P _eq是平衡压力或平台压力,ΔH ⁰和ΔS ⁰分别为标准焓变和熵变。ΔW _surf为产生新界面需要的能量;γ=(∂G/∂A)_T,P,为表面能密度,A为表面积。结合式(22)和(23),表面能ΔW _surf为负值,可降低体系的放氢焓变。因此,纳米颗粒提供的多余界面能将进一步促进纳米材料热力学上的去稳定。

目前,合成纳米级材料的方法主要有两种:(1)“自上而下”法,(2)“自下而上”法。本节就上述两种方法总结硼氢化物纳米化方面的进展。

2.4.1 高能球磨

高能球磨(自上而下)通过碰撞、研磨和搅拌来降低材料的尺寸,是一种简单又经济的纳米化方法^{[231, 232]}。然而,研磨技术仅能将颗粒尺寸降至100~200 nm,且颗粒尺寸分布宽^[49]。DFT研究表明,颗粒尺寸减至小于10 nm才可能改变体系热力学^{[233, 234]}。另一方面,由于球磨过程中颗粒间存在冷焊现象,因此需要引入防粘剂以提高球磨效率,例如球磨镁时加入的石墨碳,但防粘剂的引入牺牲了整体储氢容量^[235]。即使抛开上述问题,也很难单纯通过球磨来制备纳米级硼氢化物。研究发现,Mg(BH₄)₂球磨后尺寸约为10 μm,且颗粒易团聚。虽然球磨过程会使金属硼氢化物产生晶格缺陷,减少晶粒尺寸,但效果极其有限,其放氢温度几乎无变化^[236],仅在Ca(BH₄)₂体系中观察到10 ℃的降低^[62]。球磨并不能有效地减小金属硼氢化物的颗粒尺寸,这是因为其内在结构稳定和元素组成复杂,导致球磨对硼氢化物的吸氢性能几乎没有影响^[51]。因此,如何制备纳米金属硼氢化物,且保证其在吸放氢循环期间的稳定性,仍面临着巨大挑战。

2.4.2 沉积

Chen等利用气相输送的方法从大块Mg中合成了直径为30~50 nm的Mg纳米线。在此方法中,商业化的Mg块先被加热到蒸发温度(1203 K),然后Mg蒸汽在温度相对较低的区域(573 K)发生沉积,制备了Mg纳米线。合成的Mg纳米线几何尺寸较块状Mg大大降低,尺寸效应的存在,使体系的热力学和动力学性能发生改变,因此具有较低的反应能^[237]。受此启发,Pang等基于传统球磨和气相沉积,开发了机械力驱动物理气相沉积(MFPVD)方法,以此合成了LiBH₄纳米带(自上而下)。图6为以[Mg(AlH₄)₂(Et₂O)] _n为例给出的合成方法示意图。基于高能球磨运行时产生的局部和瞬时高温,在研磨罐中会形成微型气相沉积系统,随后有机配位聚合物[LiBH₄(MTBE)] _n(MTBE:甲基叔丁基醚)历经细化、蒸发,最后以带状形态沉积在罐盖上,宽度为10~40nm^[238]。MFPVD法得到的纳米LiBH₄在60 ℃时即可释放大量氢气,这是首次也是唯一关于直接制备纳米级硼氢化物的报道。

2.4.3 从分子组装纳米材料

球磨适用于传统金属氢化物的纳米化,但对于配位氢化物,特别是硼氢化物有诸多限制,容易导致如尺寸减少有限,粒度分布宽和颗粒易团聚等许多问题,而MFPVD的纳米化效率较低,难以控制。在无机材料领域,液相合成法,包括共沉淀、溶胶-凝胶加工、微乳液、水热/溶剂法、模板合成和仿生合成等,均可有效制备相应的纳米颗粒^[239]。通过溶胶-凝胶法,研究者成功合成了5 nm的胶体Mg,其在室温条件即能可逆储存~7.6 wt%的

H 2

^[240]。虽然液相法能基于前驱体有效制备相应的金属或金属氧化物纳米颗粒,但由于缺乏合适的反应前驱体,此方法难以应用于硼氢化物。

最近,研究者通过反溶剂沉淀法制备出小于30 nm的NaBH₄纳米颗粒(自下而上)。其中四丁基溴化铵(TBAB)表面活性剂起到了关键作用,其发挥了空间排斥作用,成功限制了纳米NaBH₄的生长。借此方法,研究者合成了NaBH₄@Ni、 NaBH₄@Co、NaBH₄@Cu、NaBH₄@Fe和NaBH₄@Sn核壳结构体系。能量色散X射线光谱(EDS)分析显示,涂层金属多存在于颗粒的外缘,而Na的浓度朝中心逐渐增加,证实了核壳结构的生成。此外,纳米NaBH₄的放氢起始温度较大块NaBH₄降低了90 ℃^{[241, 242]}。Wang等使用溶剂蒸发法及不同的表面活性剂合成了不同尺寸LiBH₄纳米粒子,表面活性剂的链长、空间位阻、活性剂端基基团与LiBH₄的结合能力等因素都会影响LiBH₄纳米颗粒的生长,表面活性剂链长越长、其与LiBH₄结合力越强,生成的LiBH₄颗粒越小。以十二胺(dodecylamine,DDA)作为表面活性剂可得到颗粒尺寸小于20 nm的LiBH₄-DDA产物,由于同时引入了—NH₂,导致NH^δ+和BH^δ-间的复合作用,协同降低了LiBH₄的放氢温度^[243]。然而,除上述研究,还没有其他金属硼氢化物以此方法合成,这主要归因于其他金属硼氢化物的液相合成缺少合适的溶剂,表面活性剂和参数控制条件。

除了制备困难外,所制纳米硼氢化物在高温下易于烧结/团聚,导致颗粒变大。由于小颗粒的高活性,目前维持纳米颗粒在吸放氢循环过程中的低维尺寸仍异常困难^[244],且其变大后难以恢复至初始尺寸^[241]。再加上高频碰撞和奥斯特瓦尔德熟化的存在^[239],小颗粒本质上易于团聚,导致原有尺寸和边界效应的消失,严重影响了储氢材料后续的吸放氢性能。

2.4.4 基于多孔框架材料的纳米限域

为了维持储氢材料的纳米尺寸,若将储氢材料限制于纳米多孔框架中,使其在既定的框架中进行吸放氢反应,有望限制其长大,这种方法被称为纳米限域。考虑到直接制备纳米硼氢化物困难,目前的大多数研究都集中在如何有效地将硼氢化物限制在多孔框架中(自下而上),利用这种间接的方式实现硼氢化物的纳米化,而纳米化的效果取决于框架的孔容、孔径和孔隙率。框架材料多为化学惰性且轻质的介孔或纳米多孔碳。借助框架的作用,纳米材料的生长被有效抑制,即使经历多次吸放氢,亦能维持稳定而优良的吸放氢性能。

目前,硼氢化物的纳米限域可分为溶剂浸渍和熔融渗透。ComaNescu等利用MTBE(甲基叔丁基醚)成功将Ca(BH₄)₂渗入至活性多孔碳材料中^[245]。但浸渍法引入的溶剂需被去除,这又增加了合成过程的复杂性,正如在2.2.2提到的,某些硼氢化物或是复合硼氢化物(LiB

H 4

^[246]、LiBH₄-NaB

H 4

^[247]、LiBH₄-KB

H 4

^[162]、LiBH₄-Mg(BH₄

) 2

^[248]、LiBH₄-Ca(BH₄

) 2

^[249])在分解前首先会熔融。熔融的硼氢化物可以通过毛细管作用直接注入至多孔框架中,氢背压还可以促进渗透过程。显然,熔融渗透的方法可以极大地简化纳米限域的过程,是一种更理想的优化硼氢化物储氢性能的方法。2008年,Gross等在280~300 ℃和氩气气氛下将LiBH₄限域在多孔碳气凝胶中,纳米限域LiBH₄的放氢动力学性能较块体LiBH₄提高了50倍^[250]。然而,并非所有硼氢化物都具有熔点或能在脱氢之前熔化,例如,Mg(BH₄)₂在整个放氢过程中都是以固态的形式,Ca(BH₄)₂的熔融也是伴随着氢气的释放^[168]。因此,选择合适的方法制备纳米尺寸硼氢化物至关重要。此外,框架的性质和硼氢化物的熔点亦对纳米限域硼氢化物的最终性能影响重大。

碳基框架是纳米限域中被使用最多的材料。将LiBH₄限域在活性碳纳米纤维中,体系的起始和主要放氢温度分别降低了128和118 ℃,且放氢副产物B₂H₆被有效抑制。三次吸放氢循环后,纳米限域LiBH₄仍能保留90%的储氢容量,且吸氢条件较大块LiBH₄明显降低^[251]。House等通过STEM和SEM,从形态和成分角度系统地表征了纳米多孔碳中LiBH₄的放氢过程。样品升温至200 ℃(采用TEM加热),形貌发生了明显变化:在颗粒外表面形成了由立方体和立方形LiH纳米晶体组成的粒状外壳;在250 ℃退火后,仍能在框架颗粒的外部和碳膜上清楚地看到LiH纳米晶体,这表明在放氢过程中LiH会从框架中渗出。为了验证从框架中渗出LiH与加热环境无关,采用DSC加热的样品亦进行了相同的TEM表征。虽然采用DSC加热的样品形貌不同,但同样观察到了LiH的渗出。因此,纳米限域LiBH₄可逆储氢容量的衰减,与储氢材料的损失密切相关^[252]。

进一步地,基于密度泛函理论计算,研究者发现LiBH₄纳米颗粒分解反应的热力学性能与反应物/产物和多孔框架的相互作用有关,而反应焓和(LiBH₄) _n团簇大小密切相关。n≤12时,团簇在石墨存在的条件下,在较低温度下就能分解,生成嵌入式的Li,而大(LiBH₄) _n团簇则生成(LiH

) n

^[253]。此外,碳基质和LiBH₄之间存在化学相互作用直至基质的缺陷饱和,上述作用使LiBH₄进一步去稳^[254]。

为了进一步改善纳米限域材料的储氢性能,研究者对框架进行了催化修饰^[255]。CoB和Ni对LiBH₄储氢性能的改善有积极作用,因此,研究者通过CoCl₂·6H₂O、NiCl₂·6H₂O和NaBH₄在溶剂中反应,制备了碳气凝胶负载的CoNiB纳米颗粒。随后将LiBH₄限制于上述催化修饰框架中,体系表观活化能(E _a)从59 kJ/mol(纯LiBH₄)急剧降低至46 kJ/mol,这源于纳米限域和催化的协同效应^[256]。同样,Shao等合成了NbF₅改性的介孔碳,在140 bar氢压下将LiBH₄加压熔渗至改性框架,得到了新型纳米限域体系LiBH₄@MC-NbF₅。较LiBH₄@MC,通过NbF₅的引入,LiBH₄的吸氢性能进一步提高,在温和条件下(200 ℃和60 bar氢压)能可逆存储约10.65 wt%

H 2

^[257]。此外,将LiBH₄限制于CeF₃修饰的活性炭中,亦可获得优良的可逆性^[258]。

除了上述碳基框架,研究者还发现其他材料的框架能够与储氢材料发生协同效应。高度有序的金属有机框架(MOFs)亦能够作为储氢材料的限域模板。将LiBH₄限制于Cu-MOF中,体系放氢温度进一步降低。机理分析显示,循环过程中,LiBH₄和Cu-MOF中的metal-O单元发生氧化还原反应,改善了体系的储氢性能^[259]。此外,研究者借助THF,将LiBH₄浸渍至介孔TiO₂中,体系峰值放氢温度较块状LiBH₄降低了90 ℃,且放氢速率加快,这说明较惰性碳基框架,非惰性多孔支架亦能作为硼氢化物的纳米限域材料^[260]。进一步地,ZnO/ZnCo₂O₄或NiMnO₃多孔颗粒纳米限域的LiBH₄体系具有更低的起始放氢和主要放氢温度,较原始LiBH₄,其分别降低了110~130 ℃和120~150 ℃,这归因于纳米限域和活性纳米粒子的协同作用^{[261, 262]}。

研究者在310 ℃和60 bar H₂压力下通过直接熔融渗透的方法,成功制备了2LiBH₄-MgH₂限制于碳气凝胶中的储氢体系。等温放氢结果显示,纳米限域2LiBH₄-MgH₂在T=425 ℃和3.4 bar 氢压下保温90 min即可释放体系90%的储氢容量,显著优于未限域体系(34%)。此外,体系在四次吸放氢循环后,仍能保持10.8 wt% H₂,相当于具备95%的可逆性^[262]。然而,以介孔碳CMK-3为框架, 2LiBH₄-MgH₂的储氢容量却随循环次数损失明显,这是由于体系组元协同作用弱,各组元仍为单独放氢^[263]。进一步地,研究者在纳米限域2LiBH₄-MgH₂体系中加入催化剂ZrCl₄,其放氢温度明显降低,动力学性能提高,可逆性增至97%^[264]。除了二元纳米限域体系,三元氢化物体系LiBH₄-MgH₂-NaAlH₄最近也成功限制于碳气凝胶框架中,动力学性能和操作温度均明显改善,且体系放氢过程简化为单一步骤^[265]。

虽然单一Mg(BH₄)₂和Ca(BH₄)₂体系的放氢温度和可逆性都优于LiBH₄,但由于LiBH₄的熔点较低(300 ℃),易于熔渗入碳基框架,所以金属硼氢化物纳米限域研究仍主要集中在LiBH₄。如2.2.2节提到,两种或多种硼氢化物复合体系在低于其单一组分熔点的温度下即可转变为液态。利用硼氢化物间存在共晶熔点的特性,研究者成功将多种复合硼氢化物体系限制于多孔框架中。0.55LiBH₄-0.45Mg(BH₄)₂在190 ℃即可熔融渗透于高表面积的碳气凝胶中。多次循环,纳米限域体系具有58%的可逆性,显著高于未限域体系(22%)^[248]。Karger等系统研究了LiBH₄-Mg(BH₄)₂限域于介孔碳载体样品在升温过程中的反应路径。结果表明,纳米限域加速了LiBH₄-Mg(BH₄)₂的放氢速率,并通过固态核磁共振证实了其反应途径得到了简化,有效抑制了[B₁₂H₁₂]^2-的形成,进而促进了体系可逆性^[266]。将具有共晶熔点的LiBH₄-Ca(BH₄)₂复合体系限域于介孔碳中,亦获得了优异的储氢性能。较块状0.68LiBH₄-0.32Ca(BH₄)₂,其放氢温度从~400 ℃降低至~300 ℃。由于纳米效应和LiBH₄/Ca(BH₄)₂间的协同作用,体系在无催化剂添加的情况下即可可逆吸收50%的储氢容量^[249]。另外,研究者利用CO₂活化的框架对不同组分LiBH₄/Ca(BH₄)₂体系进行了纳米限域研究,结果发现,0.7LiBH₄-0.3Ca(BH₄)₂体系的放氢温度降低至270 ℃,框架经CO₂活化后,能够抑制体系放氢过程中硼酸盐和氧化物的生成^[267]。

除了尺寸效应,纳米限域储氢材料中纳米氢化物和框架之间的界面效应亦对材料储氢性能的改善发挥了重要作用^[268]。Lee等通过(非)原位核磁共振发现熔化温度附近LiBH₄-Ca(BH₄)₂复合体系与介孔框架之间存在相互作用^[269]。进一步地,采用¹¹B,⁷Li和¹H核磁共振光谱法,对硼氢化物和介孔框架之间离子局部环境进行了系统研究^[270]。在液态LiBH₄的帮助下,纳米限域LiBH₄-Ca(BH₄)₂体系的共晶熔点进一步降低。此外,无论是将LiBH₄-Ca(BH₄)₂限域于介孔框架还是将其与框架进行简单球磨,LiBH₄-Ca(BH₄)₂体系在熔融过程均以固溶体的形式存在,且冷却后不变,而纯LiBH₄-Ca(BH₄)₂冷却后将发生相分离。此外,固溶体的纳米限域仅发生于介孔框架,而与框架的化学成分无关。LiBH₄-Ca(BH₄)₂体系与介孔框架的作用促进了体系的快速放氢^[270]。

为了进一步提高纳米限域共晶复合体系的吸放氢性能,研究者将NbF₅和共晶0.68LiBH₄-0.32Ca(BH₄)₂复合体系一起引入至CMK-3中。由于纳米限域和活性NbB₂催化剂之间的协同效应,LiBH₄-Ca(BH₄)₂-NbF₅@CMK-3体系的放氢温度降低至250 ℃,其在500 ℃时可以释放出约13.3 wt%的氢气,而纯LiBH₄-Ca(BH₄)₂体系仅能释放出10.4 wt%的氢气^[271]。另外,通过将0.62LiBH₄-0.38NaBH₄共晶复合物熔渗到碳气凝胶框架中,体系四次循环后仍然有70%的储氢容量(4.3 wt% H₂),明显高于未限域体系(1.6 wt% H₂)^[272]。结合共晶熔融,反应氢化物复合和纳米限域,纳米限域LiBH₄-KBH₄-Mg体系的吸放氢动力学性能明显提高,且释放出4 wt%~8 wt%氢气^[273]。

2.4.5 无模板纳米化

目前纳米限域技术在其大规模施用前仍有以下问题须克服:

(1)由于引入无氢多孔材料,体系整体储氢容量显着降低(<1 wt%),换句话说,多孔框架的密度必须非常低才能保证纳米限域体系的整体储氢容量。

(2)在浸渍过程中不可避免地存在孔隙堵塞,这极大地影响了储氢材料的均匀分布,限制了储氢性能的提高。

(3)由于浸渍和熔渗过程缓慢和扩散距离长,须多次反应才能实现储氢材料的高装载量。

目前还没有一种多孔框架材料能够解决上述所有问题。为了提高负载效率和均匀性,Jeon等在镁基体系中发明了一种新型原位纳米限域方法,即:通过有机金属Mg²⁺前体和萘基锂的还原反应合成Mg@PMMA(聚合物聚(甲基丙烯酸甲酯)),在未使用多孔模板的情况下,实现了原位纳米结构和纳米限域。所制备的Mg@PMMA基复合材料在显著改善体系放氢操作温度和动力学性能的情况下,较传统纳米限域技术,保持了极高的储氢容量(4 wt%)^[40]。基于PVA(聚乙烯醇)涂覆的LiN₃前驱体,研究者通过静电纺丝制备了其纳米纤维,PVA和LiN₃在热处理过程中会发生分解,原位生成了包覆C,同时,LiN₃分解生成的N₂(制孔剂)对纳米纤维形貌进行了修饰,最终形成了多孔碳包覆的Li₃N储氢材料^[274]。受此启发,Gu等将Ca(BH₄)₂与Ti(OEt)₄混合,在热处理过程中,Ti(OEt)₄为催化剂TiO₂前驱体,且分解能够产生有机气体,利用生成气体的穿梭效应,实现了对Ca(BH₄)₂的纳米加工,成功合成了原位引入纳米TiO₂的多孔Ca(BH₄)₂基氢化物CaB₂H₇,合成工艺见图7。与致密的Ca(BH₄)₂相比,多孔CaB₂H₇-0.1TiO₂体系的吸放氢动力学和热力学性能得到明显改善。然而,体系在首次吸放氢循环后,由于缺乏约束框架,所制备的纳米结构被破坏,影响了其后续性能^[275]。

3 结论与展望

到目前为止,研究者开展了大量研究对金属硼氢化物的储氢性能进行优化。然而,现有的金属硼氢化物及其优化技术还不能完全满足储氢的应用要求。对于替代法,工作温度、可逆储氢容量及循环性都均须进一步改善。离子替代体系由于其具有快速的离子导电性,似乎更适合用作电池的新型固态电解质;基于二元复合体系,多元复合体系的开发值得进一步研究,通过成分调变,有望获得综合性能优异的储氢体系;我们认为纳米化是解决当前问题最有希望的方法,但它仍然面临两个关键问题:实际容量和循环稳定性。除了开发新型框架,无模板纳米限域方法也值得进一步研究。毫无疑问,添加催化剂是改善储氢材料性能最直接的方法,但催化活性和稳定性须进一步改进,此外催化改性更多是扮演其他改性方法的辅助角色。总的来说,没有一个单一的方法可以克服所有问题来实现金属硼氢化物的实用性,多种方法结合协调,共同优化材料的热力学和动力学性能,才能实现氢能的大规模应用。

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